Hot-rolled steel sheet and cold-rolled steel sheet and manufacturing method thereof

ABSTRACT

A steel sheet excellent in mechanical strength, workability and thermal stability and suited for use as a raw material in such fields of manufacturing automobiles, household electric appliances and machine structures and of constructing buildings, and a manufacturing method thereof. are provided. The steel sheet is a hot-rolled steel sheet of carbon steel or low-alloy steel, the main phase of which is ferrite, and is characterized in that the average ferrite crystal grain diameter D (μm) at the depth of ¼ of the sheet thickness from the steel sheet surface satisfies the relations respectively defined by the formulas (1) and (2) given below and the increase rate X (μm/min) in average ferrite crystal grain diameter at 700° C. at the depth of ¼ of the sheet thickness from the steel sheet surface and said average crystal grain diameter D (μm) satisfy the relation defined by the formula (3) given below: 
       1.2≦D≦7  formula (1) 
       D≦2.7+5000/(5+350.C+40.Mn) 2   formula (2) 
         D·X ≦0.1  formula (3) 
     wherein C and Mn represent the contents (in % by mass) of the respective elements in the steel.

TECHNICAL FIELD

The present invention relates to a hot-rolled or cold-rolled steel sheethaving ultra fine crystal grains, and to a manufacturing method thereof.More particularly, it relates to a hot-rolled or cold-rolled steel sheetexcellent in mechanical strength, workability and thermal stability andsuited for use as a raw material in such fields of manufacturingautomobiles, household electric appliances and machine structures and ofconstructing buildings, and to a manufacturing method thereof.

BACKGROUND ART

The steel sheet to be used as a raw material for structural members intransport machines, typically automobiles, and various industrialmachines is sometimes required not only to be excellent in suchmechanical characteristics as strength, workability and toughness butalso to have weldability in parts assembly and/or corrosion resistanceduring use. In order to generally improve the mechanical characteristicsof a steel sheet, it is effective to render a steel sheet with finemicrostructure. Therefore, a number of methods in order to render asteel sheet with fine microstructure have been proposed.

To sum up, the means for rendering a steel sheet with finemicrostructure as known in the prior art include (i) the high reductionrolling method, (ii) the controlled rolling method, (iii) the alloyingelement addition method, and combinations of these.

The high reduction rolling method (i) is a technique which comprisesemploying a rolling reduction of about 50% or higher to causeaccumulation of great strains in one rolling pass and then transformingaustenite grains to fine ferrite grains or recrystallizing relativelylarge ferrite grains into fine ferrite grains caused by the greatstrains. Such technique makes it possible to obtain an ultra fineferritic microstructure with a grain size of 1 to 3 μm by heating to atemperature not higher than about 1000° C. and then carrying outhigh-reduction rolling within a low temperature zone around 700° C.However, this method is difficult to realize on an industrial scale and,in addition, has a problem that since the fine ferritic microstructurereadily allows grain growth during heat treatment, and the weldedportion, upon welding, becomes softened or, upon hot-dip Zn plating, themechanical characteristics expected are lost.

The controlled rolling method (ii) is a technique comprising carryingout multi-pass rolling generally at a temperature not lower than about800° C. employing a reduction, per rolling pass, of not higher than 20to 40%, followed by cooling. Various modifications have been disclosed,for example such as the method employing a rolling temperature within anarrow range around the Ar₃ point, the method employing a shortenedpass-to-pass time in rolling, and the method causing dynamicrecrystallization of austenite while controlling the strain rate andtemperature. However, no full investigations have been made concerningthe cooling after rolling. While it is said that water-coolingimmediately after rolling is preferred, the cooling, though said to bemade immediately, is actually started after the lapse of 0.2 second or alonger period after rolling, and the cooling rate is at most about 250°C./second. Such method can reduce the ferrite crystal grain size oflow-carbon steel simple in composition only to about 5 μm. Therefore, itis impossible to improve the mechanical characteristics to satisfactorylevels.

The alloying element addition method (iii) is to render ferrite crystalgrains fine by addition of a minute amount of at least one alloyingelement capable of suppressing recrystallization or recovery ofaustenite. Such alloying elements as Nb and Ti form carbides and/orsegregate at grain boundaries to thereby prevent recovery andrecrystallization of austenite, so that the austenite grains after hotrolling are fine and the ferrite crystal grains obtained bytransformation of austenite are also fine. This alloying elementaddition method (iii) is often used in combination with the highreduction rolling method (i) and/or controlled rolling method (ii)mentioned above. This alloying element addition method (iii) is alsoeffective in suppressing ferrite grain growth during heat treatment aswell. However, this method has a problem in that it causes reductions inaustenite volume fraction although it reduces the ferrite crystal grainsize; further, the method is yet unsatisfactory in suppressing ultrafine ferrite crystal grains from growing during welding or hot-dip Znplating. Therefore, the method is applicable only to limited steelspecies. In addition, the material cost goes up due to the alloyingelement(s) to be added.

A prior art referring to these high reduction rolling method (i),controlled rolling method (ii) and alloying element addition method isdisclosed in Patent Document 1. In that document, a method is disclosedwhich comprises finishing working in one or more passes at a totalreduction of 50% or higher within 1 second in a temperature range ofAr₁+50° C. to Ar₃+100° C. and carrying out forced cooling, afterfinishing of the working, within a temperature range of not lower than600° C. at a cooling rate of not lower than 20° C./second.

Further, Patent Document 2 discloses a method which comprises carryingout rolling passes in at least 5 stands within the dynamicrecrystallization temperature range, with the temperature differencebetween the first stand entry side and the last stand outlet side beingnot greater than 60° C.

[Patent Document 1] JP-S59-205447-A

[Patent Document 2] JP-H11-152544-A

DISCLOSURE OF INVENTION Problems to be Solved

However, even when steel sheets having a fine crystal microstructure areobtained by such methods, the microstructure is low in thermalstability. Therefore, even if the microstructure is rendered fine andthe mechanical characteristics are improved, the subsequent step ofwelding of steel sheets or hot-dip plating of steel sheets readilyresults in coarsening of crystal grains due to heating in that step andthe mechanical characteristics of the sheets are seriously impaired;this is another problem. Further, when these hot-rolled steel sheets aresubjected to cold rolling and heat treatment to give thin steel sheets,a further problem arises, namely the heat treatment readily results incoarsening of crystal grains and cold-rolled steel sheets having a finemicrostructure can scarcely be obtained.

It is an objective of the present invention to provide a hot-rolled orcold-rolled steel sheet excellent in thermal stability and mechanicalcharacteristics and capable of enduring the heat during welding orhot-dip plating, and a method of producing the same.

Means for Solving the Problems

The present inventors made various investigations and experimentsconcerning the mechanical characteristics and thermal stability of thefine ferrite crystal grain microstructure and, as a result, found that,in order to attain both good mechanical characteristics and good thermalstability, it is most important (a) for the average ferrite crystalgrain diameter to be kept within a certain specific range and (b) to setan upper limit to the product D·X (μm²/min) of the increase rate X(μm/min) in average ferrite crystal grain diameter D (μm) attemperatures in the vicinity of about 700° C. just below the A₁ pointand that average crystal grain diameter D (μm). They also found that, inorder to attain better thermal stability, it is preferable (c) to retainthe ferrite crystal grain distribution within a certain specific rangeand/or for no strains resulting from rolling to be left within ferritecrystal grains Further, they found that steel sheets obtained by takingsuch findings into consideration, when cold-rolled and thenheat-treated, can again acquire a ferrite crystal grain microstructurethermally stable and fine in the same manner as mentioned above. Then,they made various investigations and experiments in search of (d) anovel method of producing hot-rolled or cold-rolled steel sheets havingsuch a microstructure and such characteristics. Furthermore, as regardswelded members, they found (e) that, in fusion welding, it is preferableto specify the hardness balance in the weld and (f) that, in resistancewelding, it is preferable to strive for attaining a hardness balance andpreventing embrittlement in the weld.

In the following, the findings and results of investigations andexperiments with regard to the above-mentioned items (a) to (f), whichhave led to completion of the present invention, are described indetail.

(a) Re: To Keep the Average Ferrite Crystal Grain Diameter within aSpecific Range:

As the ferrite crystal grain diameter increases, the strength increases.It was found, however, that when the crystal grains become too small,the driving force for grain growth owing to grain boundary energybecomes increased and, as a result, the grain growth at hightemperatures is promoted. More specifically, it was revealed that whenthe average crystal grain diameter is smaller than 1.2 μm, it becomesdifficult to control the grain growth at high temperatures and,conversely, when the average crystal grain diameter is in excess ofeither the value of 2.7+5000/(5+350.C+40.Mn)² μm or the value of 7 μm inthe case of hot-rolled steel sheets or, in the case of cold-rolled steelsheets, in excess of either the value of5.0−2.0.Cr+5000/(5+350.C+40.Mn)² μm or the value of 9.3 μm, satisfactoryimprovements in mechanical characteristics as a result of rendering themicrostructure fine can no longer be expected. Therefore, in order toattain both satisfactory mechanical characteristics and good thermalstability simultaneously, it is necessary to employ the value of 1.2 μmas the lower limit to the average ferrite crystal grain diameter and, asthe upper limit thereto, the value smaller of the value of2.7+5000/(5+350.C+40.Mn)² μm and the value of 7 μm in the case ofhot-rolled steel sheets or, in the case of cold-rolled steel sheets, thevalue smaller of the value of 5.0−2.0.Cr+5000/(5+350.C+40.Mn)² μm andthe value of 9.3 μm.

(b) Re: To Set an Upper Limit to the Product D·X of the Increase Rate Xin Average Ferrite Crystal Grain Diameter D at Temperatures in theVicinity of about 700° C. Just Below the A₁ Point:

The rate of ferrite crystal grain growth at high temperatures increaseswith the increase in temperature. Generally, the temperature range inwhich the ferrite grain growth problem arises during welding or hot-dipplating is a temperature range from just below the A₁ point (about 730°C.) to the vicinity of the A₃ point and, within this temperature range,the rate of ferrite grain growth varies markedly. However, since it wasrevealed that the temperature characteristics of the grain growth ratein steel sheets of which the average ferrite crystal grain diameter iswithin the range specified at the item (a) above are determined by therate of ferrite grain growth at temperatures in the vicinity of 700° C.,it was found that when an upper limit is set to the rate of ferritegrain growth at temperatures in the vicinity of 700° C., namely to theproduct D·X (μm²/min) of the increase rate X (μm/min) in average ferritecrystal grain diameter D (μm) and the average crystal grain diameter D(μm), no problems will occur even when heating is made at highertemperatures during welding or hot-dip plating. And, as a result ofexperiments, it was revealed that it is necessary to set the product D·Xat 0.1 μm²/min or below. The product D·X is preferably not higher than0.07 μm²/min, and more preferably not higher than 0.05 μm²/min.

(c1) Re: To Retain the Ferrite Crystal Grain Distribution within aSpecific Range and to Leave None of the Strains Resulting from Rollingwithin Ferrite Crystal Grains:

The ferrite crystal grain diameter distribution and the strains withinferrite crystal grains are closely associated with the grain growth athigh temperatures. The grain growth at high temperatures occurs with thegrain boundary energy and strains within grains serving as drivingforces. Therefore, when relatively large ferrite crystal grains coexistin a fine ferritic microstructure, the large ferrite crystal grainsreadily integrate with surrounding fine ferrite crystal grains utilizingthe boundary as a driving force. When there are strains within ferritecrystal grains, the intragranular strains serve as driving forces toreadily unify neighboring ferrite crystal grains with each other. Inthis manner, the grain growth proceeds rapidly. Therefore, in order toprevent the rapid progress of grain growth, it is preferred that, inaddition to rendering ferrite crystal grains fine, the ferrite crystalgrain diameter distribution be such that at least 80% of the grains havea diameter falling within the range of ⅓ to 3 times the average crystalgrain diameter. This crystal grain diameter distribution is measured ata specified depth from the sheet surface or within the range of 100 μmfrom that depth. This is because, while the crystal grain diameter inthe steel sheet obtained according to the method of the invention variesin the direction of sheet thickness, as described later herein, the mildchanges in crystal grain diameter in such direction of sheet thicknessdoes not influence the growing behavior of grains. The intragranulardislocation density indicative of the strains within ferrite crystalgrains is preferably not higher than 10⁹/cm², and more preferably nothigher than 10⁸/cm². Further, the ferrite grains are preferably equiaxedin shape.

(c2) Re: Ferrite Grain Diameter Distribution in the Direction of SheetThickness:

A gentle ferrite grain diameter distribution in the direction of sheetthickness which shows a tendency for the grains to become finer from thecentral portion of the steel sheet to the steel sheet surface layer ispreferred from the viewpoint of improvements in steel sheet workability,for example in hole expandability or bendability. The ferriticmicrostructure rendered fine in the surface layer also contributes toimprovements in steel sheet surface treatability, for example inchemical treatability or platability. Therefore, in the case ofhot-rolled steel sheets, it is preferred that the average crystal graindiameter Ds (μm) at the depth of 1/16 of the sheet thickness from thesteel sheet surface, the average crystal grain diameter D (μm) at thedepth of ¼ of the sheet thickness from the steel sheet surface and theaverage crystal grain diameter Dc (μm) in the central portion of thesheet thickness satisfy the relations Ds≦0.75Dc and D≦0.9Dc and, in thecase of cold-rolled steel sheets, it is preferred that the relationDs≦0.9Dc be satisfied.

(d) Re: Novel Method in Order to Produce Hot-Rolled Steel Sheets Havingthe Microstructure and Characteristics Mentioned in the Items (a) to (c)Above:

By employing the method of rolling in a high temperature range asdescribed below, it becomes possible to provide an easy andhigh-productivity industrial method of rolling.

First, multi-pass hot rolling is started from the austenite temperaturerange, and the final rolling pass is finished at a high temperature notlower than the Ar₃ point and not lower than 780° C. On that occasion,strains are accumulated within austenite crystal grains.

And, within 0.4 second directly after the finish of hot rolling, coolingto a temperature not higher than 720° C. is completed. On that occasion,the strains are suppressed from being relieved during cooling and thestrains remain accumulated within austenite grains and, at temperaturesnot higher than 720° C., the transformation of austenite to ferritefirst becomes active and a large number of ferrite crystal grains aregenerated with the accumulated strains serving as nuclei, forming a fineferritic microstructure. According to this method, the shear strainsintroduced into the steel sheet during hot rolling as a result offriction between the steel sheet surface and the rolling roll surfacecan also be suppressed from being relieved, so that a larger number offerrite nuclei are generated in the portion closer to the surface thanin the sheet thickness center portion.

Further, thereafter, the steel sheet is maintained in a temperaturerange of 600 to 720° C. for at least 2 seconds. By this measure, itbecomes possible to obtain a desired ferritic microstructure withcrystal grains fine in diameter being distributed in a narrow range and,at the same time, the strains are suppressed from remaining in the fineferritic microstructure after transformation. Further, theabove-mentioned changes in the amount of ferrite nuclei formed in thedirection of sheet thickness contribute to the formation of amicrostructure showing a moderate grain diameter gradient from the sheetthickness center to the surface.

As for the cooling conditions just after completion of the hot rolling,it is necessary that the cooling to a temperature of 720° C. or below befinished within 0.4 second, as described above. In the prior art,cooling is started after the lapse of at least 0.2 second, at thespeediest, just after completion of the rolling, and the rate of coolingwas at most about 250° C./second. In the case of a low-carbon steelspecies whose Ar₃ point is 800° C., for instance, the prior art requiresat least 0.52 second for cooling from 800° C. or above to a temperaturenot higher than 720° C. even when the hot rolling of the low-carbonsteel is finished at the Ar₃ point; thus, in the art, it is difficult tofinish the cooling to 720° C. or below within 0.4 second.

When a hot-rolled steel sheet having such a microstructure as describedin the items (a) to (c) above is cold rolled and then heat-treated at atemperature not higher than the temperature (Ac₃) at which an austeniticsingle phase is formed, a fine-grained ferritic microstructure havingthe characteristics mentioned above is again obtained. This ispresumably due to (1) a generation of a large number of ferrite nucleicaused by a generation, on abundantly existing ferrite grain boundariesresulting from hot rolling, of nuclei for recrystallization of processedferrite during heat treatment after cold rolling and (2) simultaneousgeneration of a large number of austenite grains on ferrite grainboundaries to suppress the growth of ferrite nuclei. As a result, theferrite grain diameter is almost the same as or larger only by 1 to 3 μmthan the ferrite diameter obtained upon hot rolling and, thus, amicrostructure inheriting the characteristics acquired on the occasionof hot rolling is obtained. Therefore, when a ferrite crystal graindiameter distribution in the direction of sheet thickness is found inthe stage of hot-rolled steel sheets, as in accordance with the presentinvention, a good ferrite crystal grain diameter distribution in thedirection of sheet thickness can appear even after cold rolling and heattreatment. While the heat treatment temperature may be at a level of Ac₁or below, a long period of time is required for the recrystallization ofprocessed ferrite. At a temperature not lower than Ac₁ at which anaustenitic single phase appears, the microstructure tends to becomecoarsened with ease.

(e) Re: To Specify the Hardness Balance at the Weld in Fusion Welding:

In arc welding in which the heat input on the occasion of welding islarge, it goes without saying that, from the viewpoint of preventing theheat-affected zone (HAZ) from being softened, it is preferable to form ahighly thermally stable microstructure scarcely causing grain growthduring welding. Furthermore, in order to secure the workability of amember after welding, it is preferable to specify the hardness balanceof the weld and thereby improve the fusion weldability. Thus, as regardsthe chemical composition, it is possible to obtain a weldable partexcellent in fusion weldability by prescribing that the carbonequivalent Ceq(I), which is defined byCeq(I)=C+Mn/6+Si/24+Cr/5+Mo/4+Ni/40+V/14, should be 0.06 to 0.6%. Thefusion weldability refers to such characteristic that the differencebetween the maximum hardness of the weld obtained by using a weldingmethod which proceeds with continuous molten pool forming and moltenpool solidification, such as in arc welding or laser beam welding, andthe hardness of the base metal or the hardness of the most softened partof the weld is small and that the weld is prevented from beingembrittled, hence the workability of the member after welding can besecured.

(f) Re: To Strive for Attaining a Hardness Balance and Preventing theWeld from being Embrittled in Resistance Welding:

In resistance welding as well, in which welding is effected byelectrothermal heating of the base metal, it goes without saying that itis preferable to form a highly thermally stable microstructure scarcelycausing grain growth during welding. Furthermore, it is preferable tostrive for attaining a hardness balance and preventing embrittlement inthe weld. Thus, as regards the chemical composition, it is possible toobtain a weldable part excellent in resistance weldability byprescribing that the carbon content should be C≦0.17 and that the carbonequivalent Ceq(II), which is defined by Ceq(II)=C+Mn/100+Si/90+Cr/100,should be 0.03 to 0.20% and, further, that the indicator Rsp of the basemetal resistance, which is defined by Rsp=13.5×(Si+Al+0.4Mn+0.4Cr)+12.2,should be not higher than 45 so that sufficient weld nuggets forsecuring the joint strength may be obtained within a broad range ofwelding conditions. The resistance weldability refers to suchcharacteristic that a sufficient level of joint strength (the so-calledmaximum breaking force at button breakage) can be secured within a broadrange of welding conditions.

The present invention has been completed based on such findings andinvestigational and experimental results. The present invention consistsin a hot-rolled steel sheet as defined below in any of the items (1),(2), (4) to (7) and (9) to (11), a cold-rolled steel sheet as definedbelow in any of items (3) to (6) and items (8) to (11) and, further, amethod of hot-rolled steel sheet production as defined below under anyof items (12) and (14) and a method of cold-rolled steel sheetproduction as defined below under any of items (13) and (14).Hereinafter, the inventions defined under those items are respectivelyreferred to as the inventions (1) to (14). The inventions (1) to (14)are sometimes collectively referred to as the present invention.

The carbon steel or low-alloy steel to be used in the practice of theinvention preferably contains 0.01 to 0.25% of C and may further containone or more elements selected from among Si, Mn, Al, P, Ti, Nb, V, Cr,Cu, Mo, Ni, Ca, REM and B.

Item (1)

A hot-rolled steel sheet of carbon steel or low-alloy steel, the mainphase of which is ferrite, and is characterized in that the averageferrite crystal grain diameter D (μm) at the depth of ¼ of the sheetthickness from the steel sheet surface satisfies the relationsrespectively defined by the formulas (1) and (2) given below and, at thesame time, the increase rate X (μm/min) in average ferrite crystal graindiameter at 700° C. at the depth of ¼ of the sheet thickness from thesteel sheet surface and the above-mentioned average crystal graindiameter D (μm) satisfy the relation defined by the formula (3) givenbelow:

1.2≦D≦7  formula (1)

D≦2.7+5000/(5+350.C+40.Mn)²  formula (2)

D·X≦0.1  formula (3)

Here, C and Mn indicate the contents (in % by mass) of the respectiveelements in the steel.

Item (2)

A hot-rolled steel sheet according to item (1) above, characterized inthat, at the depth of ¼ of the sheet thickness from the steel sheetsurface, the area percentage of ferrite crystal grains the crystal graindiameter d (μm) of which satisfies the relation defined by the formula(4) given below amounts to at least 80%:

D/3≦d≦3D  formula (4)

Here, D represents the average ferrite crystal grain diameter (μm) atthe depth of ¼ of the sheet thickness from the steel sheet surface.

Item (3)

A cold-rolled steel sheet of carbon steel or low-alloy steel, the mainphase of which is ferrite, and is characterized in that the averageferrite crystal grain diameter D (μm) at the depth of ¼ of the sheetthickness from the steel sheet surface satisfies the relationsrespectively defined by the formulas (5) and (6) given below and, at thesame time, the increase rate X (μm/min) in average ferrite crystal graindiameter at 700° C. at the depth of ¼ of the sheet thickness from thesteel sheet surface and the above-mentioned average crystal graindiameter D (μm) satisfy the relation defined by the formula (3) givenbelow:

1.2≦D≦9.3  formula (5)

D≦5.0−2.0.Cr+5000/(5+350.C+40.Mn)²  formula (6)

D·X≦0.1  formula (3)

and, further, that, at the depth of ¼ of the sheet thickness from thesteel sheet surface, the area percentage of ferrite crystal grains thecrystal grain diameter d (μm) of which satisfies the relation defined bythe formula (4) given below amounts to at least 80%:

D/3≦d≦3D  formula (4)

Here, C, Cr and Mn represent the contents (in % by mass) of therepresent element in the steel.

Item (4)

A hot-rolled or cold-rolled steel sheet according to any of items (1) to(3) above, characterized in that it contains, as a second phase otherthan ferrite, a total of less than 50% of one or more species selectedfrom the group consisting of less than 50% of bainite, less than 30% ofpearlite, less than 5% of granular cementite, less than 5% of martensiteand less than 3% of retained austenite and has a yield ratio lower than0.75, % in each occurrence being % by volume.

Item (5)

A hot-rolled or cold-rolled steel sheet according to any of items (1) to(3) above, characterized in that it contains, as a second phase otherthan ferrite, 5 to 40% by volume of martensite and has a yield ratiolower than 0.75.

Item (6)

A hot-rolled or cold-rolled steel sheet according to any of items (1) to(3) above, characterized in that it contains, as a second phase otherthan ferrite, 3 to 30% by volume of retained austenite and has aproduct, TS×El, of tensile strength TS (MPa) and total elongation El (%)of not less than 18000 (MPa·%).

Item (7)

A hot-rolled steel sheet according to any of items (1), (2), (4), (5)and (6) above, characterized in that the average crystal grain diameterDs (μm) at the depth of 1/16 of the sheet thickness from the steel sheetsurface, the average crystal grain diameter D (μm) at the depth of ¼ ofthe sheet thickness from the steel sheet surface and the average crystalgrain diameter Dc (μm) at the center of the sheet thickness satisfy therelations Ds≦0.75Dc and D≦0.9Dc.

Item (8)

A cold-rolled steel sheet according to any of items (3) to (6) above,characterized in that the average crystal grain diameter Ds (μm) at thedepth of 1/16 of the sheet thickness from the steel sheet surface andthe average crystal grain diameter Dc (μm) at the center of the sheetthickness satisfy the relation D≦0.9Dc.

Item (9)

A hot-rolled or cold-rolled steel sheet according to any of items (1) to(8) above, characterized in that the carbon equivalent Ceq(I) defined bythe formula (7) given below is 0.06 to 0.6%:

Ceq(I)=C+Mn/6+Si/24+Cr/5+Mo/4+Ni/40+V/14  formula (7)

Here, the symbols of elements in the above formula represent thecontents (in % by mass) of the respective elements in the steel.

Item (10)

A hot-rolled or cold-rolled steel sheet according to any of items (1) to(8) above, characterized in that the C content is not higher than 0.17%by mass, the carbon equivalent Ceq(II) defined by the formula (8) givenbelow is 0.03 to 0.20% and the base metal resistance indicator Rspdefined by the formula (9) given below is not higher than 45:

Ceq(II)=C+Mn/100+Si/90+Cr/100  formula (8)

Rsp=13.5×(Si+Al+0.4Mn+0.4Cr)+12.2  formula (9)

Here, the symbols of elements in the above formulas represent thecontents (in % by mass) of the respective elements in the steel.

Item (11)

A hot-dip-plated hot-rolled or cold-rolled steel sheet characterized inthat it comprises a Zn, Al, Zn—Al alloy or Fe—Zn alloy coat layer formedon the surface of a hot-rolled steel sheet according to any of items (1)to (10) above.

Item (12)

A method of producing a hot-rolled steel sheet according to any of items(1), (2), (4), (5), (6), (7), (9), (10) and (11) above by subjecting acarbon steel or low-alloy steel slab to multi-pass hot rolling to givethe hot-rolled steel sheet, characterized in that the final rolling passis finished at a temperature not lower than the Ar₃ point and not lowerthan 780° C. and then the rolled sheet is cooled to 720° C. or belowwithin 0.4 second at a cooling rate of not lower than 400° C./second andthen maintained in a temperature range of 600 to 720° C. for at least 2seconds.

Item (13)

A method of producing a cold-rolled steel sheet according to any ofitems (3), (4), (5), (6), (8), (9) and (10) above, characterized in thata hot-rolled steel sheet obtained by the method according to item (12)above is pickled, then cold-rolled at a reduction of 40 to 90% andthereafter heat-treated at a temperature of not higher than 900° C.

Item (14)

A method of producing a hot-dip-plated hot-rolled or cold-rolled steelsheet according to item (11) above, characterized in that a hot-rolledsteel sheet obtained by the method according to item (12) above issubjected to pickling or to picking and further cold rolling at areduction of 40 to 90%, and then to hot-dip plating in a continuoushot-dip plating line.

EFFECTS OF THE INVENTION

According to the present invention, a hot-rolled steel sheet and acold-rolled steel sheet, each having ultra fine crystal grains, havingsuch thermal stability that allows the steel sheet to endure heat duringwelding or hot-dip plating and, further, excellent in mechanicalcharacteristics as well as a manufacturing method thereof can beprovided.

BEST MODES FOR CARRYING OUT THE INVENTION

In the following, the ultra fine crystal grain steel sheet according tothe present invention is described. In the following, “%” used inrelation to the content of each chemical component means “% by mass”.

(A) Re: Chemical Composition:

C:

C is an element useful in promoting the process of rendering ferritecrystal grains fine since it can lower the austenite-to-ferritetransformation temperature and lower the finish temperature in hotrolling. It is also an element in order to secure the strength.Therefore, it is preferred that not lower than 0.01% of C be contained.In order to promote the process of rendering ferrite crystal grainsfine, the content is preferably not lower than 0.03%. Since, however, anexcessive content may result in retardation of ferrite transformationafter hot rolling and in decreases in volume fraction of ferrite and,further, in deterioration in weldability, the content is preferably nothigher than 0.25%. In order to improve the workability of the weldablepart, the C content is preferably not higher than 0.17%, and morepreferably not higher than 0.15%.

Si:

Si is preferably contained in order to improve the strength. However, anexcessive content results in marked deterioration in ductility, causinga surface oxidation during hot rolling. Therefore, it is preferable thatthe content be not higher than 3%, more preferably not higher than 2%,and much more preferably not higher than 1.8%. The lower limit may be animpurity level. In order to form a retained austenite in the ferritemicrostructure, the total content of Si+Al is preferably not lower than1%.

Mn:

Mn is preferably contained in order to secure the desired strength.Further, it makes it possible to lower the austenite-to-ferritetransformation temperature and lower the finish temperature in hotrolling and thus promote the process of rendering ferrite crystal grainsfine and, therefore, it is preferable that it be contained in the steel.However, an excessive content causes retardation of the ferritetransformation after hot rolling and lower the volume fraction offerrite and, therefore, the content is preferably not higher than 3%,more preferably not higher than 2.7%. The lower level may be an impuritylevel but, in the case of addition for strength improvement, the contentis preferably not lower than 0.5%. In order to cause a retainedaustenite to be formed in the ferrite microstructure, the content ispreferably not lower than 0.5%, and more preferably not lower than 0.8%.In order to form martensite in the ferrite microstructure, the contentis preferably not lower than 1.5%.

Al:

Al can be added for ductility improvement. However, an excessive contentrenders austenite unstable, making it necessary to excessively increasethe finish temperature in hot rolling and, further, making it difficultto conduct continuous casting stably. Therefore, the content ispreferably not higher than 3%. The lower limit may be an impurity levelbut, in order to cause a retained austenite formation in the ferritemicrostructure, the content is preferably such that the total content ofSi+Al is not lower than 1%.

P;

P can be added to increases the strength. However, an excessive contentcauses an embrittlement due to grain boundary segregation and therefore,when it is added, the content is preferably not higher than 0.5%, morepreferably not higher than 0.2%, and much more preferably not higherthan 0.1%. The lower limit may be an impurity level. Generally, however,an amount of about 0.01% may be intermixed into the steel during steelmaking process.

Ti;

Ti precipitates out as a carbide or nitride which increases the strengthand, further, this precipitate prevents austenite or ferrite frombecoming coarse, thus promotes the process of rendering crystal grainsfine during hot rolling and prevents grain growth during heat treatment;therefore, it can be added. However, an excessive content causes theformation of a large amount of coarse Ti carbide or nitride duringheating prior to hot rolling, which impairs the ductility and/orworkability; therefore, the content is preferably not higher than 0.3%.In order to facilitate the formation of ferrite, it is preferably addedin an amount such that the total content of Ti+Nb is not higher than0.1%, more preferably not higher than 0.03%, and more preferably nothigher than 0.01%. The lower limit may be an impurity level but it maybe generally intermixed into the steel in an amount of about 0.001%during steel making process.

Nb:

Nb precipitates out as a carbide or nitride which increases the strengthand, further, this precipitate prevents austenite or ferrite frombecoming coarse, thus promotes the process of rendering crystal grainsfine during hot rolling and prevents grain growth during heat treatment;therefore, it can be added. However, an excessive content causes theformation of a large amount of coarse NbC during heating prior to hotrolling, which impairs the ductility and/or workability; therefore, thecontent is preferably not higher than 0.1%. In order to facilitate theformation of ferrite, it is preferably added in an amount such that thetotal content of Ti+Nb is not higher than 0.1%, more preferably nothigher than 0.03%, and much more preferably not higher than 0.01%. Thelower limit may be an impurity level. Generally, however, it may beintermixed into the steel in an amount of about 0.001% during steelmaking process.

V:

V precipitates out as a carbide and increases the strength and, further,this precipitate prevents ferrite from becoming coarse and thus promotesthe process of rendering crystal grains fine; therefore, it can beadded. However, for the same reasons as in the case of Ti or Nb, itimpairs the ductility and/or workability and, therefore, the content ispreferably not higher than 1%, more preferably not higher than 0.5%, andmuch more preferably not higher than 0.3%. The lower limit may be animpurity level. Generally, however, it may be intermixed into the steelin an amount of about 0.001% during steel making process.

Cr:

Cr is effective in increasing the hardenability and causing theformation of martensite or bainite in the ferrite microstructure and,therefore, it can be added for obtaining such effects. However, anexcessive content prevents the formation of ferrite and, therefore, thecontent is preferably not higher than 1%. The lower limit may be animpurity level. Generally, however, it may be intermixed into the steelin an amount of about 0.02% during steel making process.

Cu:

Cu precipitates out at low temperatures and shows an effect ofincreasing the strength; hence it can be added for producing theseeffects. However, it may possibly cause grain boundary cracking of theslab and, therefore, the content is preferably not higher than 3%, andmore preferably not higher than 2%. When it is added, the content ispreferably not lower than 0.1%. The lower limit may be an impuritylevel. Generally, however, it may be intermixed into the steel in anamount of about 0.02% during steel making process.

Ni:

Ni can be added in order to increase the stability of austenite at hightemperatures. Further, when the steel also contains Cu, Ni can be addedin order to prevent a grain boundary embrittlement of the slab. However,an excessive content suppresses the formation of ferrite, so that thecontent is preferably not higher than 1%. The lower limit may be animpurity level. Generally, however, it may be intermixed into the steelin an amount of about 0.02% during steel making process.

Mo:

Mo precipitates out as MoC and increases the strength and, further, thisprecipitate suppresses the coarsening of ferrite and promotes theprocess of rendering crystal grains fine; therefore, it can be added.However, for the same reasons as in the case of Ti and Nb, it impairsthe ductility and workability and, therefore, the content is preferablynot higher than 1%, more preferably not higher than 0.5%, and much morepreferably not higher than 0.3%. The lower limit may be an impuritylevel. Generally, however, it may be intermixed into the steel in anamount of about 0.001% during steel making process.

Ca, REM, B:

Ca, rare earth metals (REM) or B renders finer oxides or nitridesprecipitated during solidification and maintain the soundness of theslab; therefore, one or more of them can be added. However, they areexpensive, so that the total content is preferably not higher than0.005%. The lower limit may be an impurity level.

Other “impurities” which may be intermixed into the steel are S, N andSn. As regards S and N, it is desirable that the contents be controlledin the following manner, if possible.

S:

S is an impurity element that forms a sulfide inclusion and lowers theworkability, and, therefore, it is desirable that the content be nothigher than 0.05%. When it is desired that a much higher level ofworkability be secured, the content of S is preferably not higher than0.008%, and more preferably not higher than 0.003%.

N:

N is an impurity element that decreases the workability, and it isdesirable that the content be suppressed to 0.01% or lower, and morepreferably to 0.006% or lower.

(B) Re: Microstructure of the Steel Sheet According to the PresentInvention

The steel sheet according to the present invention is a steel sheetcomprising the main phase, which is ferrite, and a second phase otherthan ferrite. The “main phase” so referred to herein means the phaseforming the largest proportion of the microstructure among the phasesconstituting the microstructure. The volume fraction of the main phaseferrite is preferably at least 50%, and more preferably not smaller than60%. When the volume fraction of ferrite is smaller than 50%, theductility and/or workability of the steel sheet may be impaired in someinstances.

The ferrite crystal grain size (diameter) exerts great influences on themechanical characteristics and thermal stability and, further,workability of the steel sheet.

Therefore, in order to secure sufficient levels of strength, ductility,thermal stability and, further, workability for the hot-rolled steelsheet according to the present invention, it is necessary that theaverage ferrite crystal grain diameter D (μm) at the depth of ¼ of thesheet thickness from the steel sheet surface be within a specific rangewithin which the relations defined by the following formulas (1) and (2)are satisfied:

1.2≦D≦7  formula (1)

D≦2.7+5000/(5+350.C+40.Mn)²  formula (2)

Thus, the specific range is a range such that the lower limit is 1.2 μmand the upper limit is the smaller value of 2.7+5000/(5+350.C+40.Mn)² μmand 7 μm.

In order to secure sufficient levels of strength, ductility, thermalstability and, further, workability for the cold-rolled steel sheetaccording to the present invention, it is necessary that the averageferrite crystal grain diameter D (μm) at the depth of ¼ of the sheetthickness from the steel sheet surface be within a specific range withinwhich the relations defined by the following formulas (5) and (6) aresatisfied:

1.2≦D≦9.3  formula (5)

D≦5.0−2.0.Cr+5000/(5+350.C+40.Mn)²  formula (6)

Thus, the specific range is a range such that the lower limit is 1.2 μmand the upper limit is the smaller value of5.0−2.0.Cr+5000/(5+350.C+40.Mn)² μm and 9.3 μm.

The reason why the lower limit to the average ferrite crystal graindiameter D is herein set at 1.2 μm is that when the diameter D issmaller than 1.2 μm, the work hardening coefficient decreases to anextreme and not only the ductility and/or workability is deteriorate butalso the thermal stability of the fine ferrite microstructure isdeteriorated and thus grains readily grow at high temperatures. In orderto obtain still higher levels of ductility, workability and thermalstability, the lower limit to the average ferrite crystal grain diameterD is preferably set at 1.5 μm. On the other hand, the reason why theupper limit to the average ferrite crystal grain diameter D is set atthe smaller value of 2.7+5000/(5+350.C+40.Mn)² μm and 7 μm forhot-rolled steel sheets or at the smaller value of5.0−2.0.Cr+5000/(5+350.C+40.Mn)² μm and 9.3 μm for cold-rolled steelsheets is that when the diameter D is in excess of any of these values,it is no more possible to obtain sufficient levels of strength. In orderto obtain still higher levels of strength, it is preferable that theupper limit to the average ferrite crystal grain diameter D be set atthe smaller value of 2.4+5000/(5+350.C+40.Mn)² μm and 5.5 μm forhot-rolled steel sheets or at the smaller value of4.5+5000/(5+350.C+40.Mn)² μm and 8.5 μm for cold-rolled steel sheets.Here, a region surrounded by large angle grain boundaries having acrystal orientation difference not smaller than 15° is defined as onecrystal grain while grain boundaries smaller in angle than 15° areneglected.

In order to further increase the thermal stability of the steel sheet,it is preferable that the ferrite crystal grain diameter distribution bewithin a specific range. A cause of the growth of grains at hightemperatures is the grain boundary energy-based driving force and, whenrelatively large ferrite crystal grains coexist in a fine ferritemicrostructure, the large ferrite crystal grains readily become unitedwith surrounding fine ferrite crystal grains with the grain boundaryserving as a driving force, whereby the grain growth progresses rapidly.Therefore, in order to reduce the rate of ferrite crystal grain growthat high temperatures, it is preferable, in addition to rendering ferritecrystal grains fine in order to keep the average crystal grain diameterD (μm) within a specific range in which the relations defined by theformulas (1) and (2) given above are satisfied, that those crystalgrains the crystal grain diameter d (μm) of which satisfies the relationdefined by the formula (4) given below amount to at least 80% in areafraction among the ferrite grains at the depth of ¼ of the sheetthickness from the steel sheet surface:

D/3≦d≦3D  formula (4)

In other words, it is preferred that the grain diameter distribution besuch that 80% or more, in area fraction, of the ferrite crystal grainsfall within the range of from one third of to 3 times the averagecrystal grain diameter D (μm). It is more preferred that the graindiameter distribution be such that 85% or more of the ferrite crystalgrains fall within the range of from ⅓ of to 3 times the average crystalgrain diameter D (μm), and much more preferably such that 90% or more ofthe ferrite crystal grains fall within the range of from ⅓ of to 3 timesthe average crystal grain diameter D (μm).

The reason why the ferrite crystal grain diameter and distribution aredefined at the depth of ¼ of the sheet thickness is that the ferritecrystal grain diameter in the hot-rolled steel sheet according to thepresent invention varies in the direction of sheet thickness. When theferrite crystal grain microstructure at that depth is maintained withinthe range mentioned above, the steel sheet according to the presentinvention can secure the desired levels of mechanical characteristicsand thermal stability. In particular, the thermal stability of the graindiameter depends not on the grain diameter distribution as statisticallydetermined in a wide range from the sheet surface to the inside but onthe grain diameter distribution as statistically determined at aspecific depth. Therefore, the statistical microstructure observationshould be carried out in a cross-section at the depth of ¼ of the sheetthickness and parallel to the surface or, if the observation is carriedout in a section perpendicular to the surface, in a region at the depthof ¼ of the sheet thickness±100 μm at the most.

The second phase other than ferrite may be any phase known to begenerally formed in low-carbon steel materials, for example pearlite,cementite, bainite, martensite, retained austenite and/or carbonitridesof elements other than Fe.

In order to efficiently produce steel sheets excellent in mechanicalcharacteristics and thermal stability, with a yield ratio of not lowerthan 0.75, it is preferred that the second phase comprises a totalamount smaller than 50%, more preferably smaller than 40%, of one ormore species selected from the group consisting of less than 50% ofbainite, less than 30% of pearlite, less than 5% of granular cementite,less than 5% of martensite and less than 3% of retained austenite, each% being percent by volume. When the volume fractions of bainite,pearlite and granular cementite are in excess of the respective valuesmentioned above, the workability is impaired. When the volume fractionsof martensite and retained austenite are in excess of the respectivevalues mentioned above, it becomes difficult to attain a yield ratio ofnot lower than 0.75.

Then, in order to efficiently produce steel sheets excellent inmechanical characteristics and thermal stability, with a yield ratio oflower than 0.75, it is preferred that the second phase comprises 5 to40% by volume of martensite. In this case, it is preferred that thevolume fractions of bainite, pearlite and granular cementite be as lowas possible. Retained austenite may be present but, in order to lowerthe yield ratio with ease, the volume fraction is preferably not higherthan 3%.

In order to efficiently produce steel sheets excellent in elongationcharacteristics, in particular, with the product of tensile strength TSand total elongation El being not smaller than 18000, and also excellentin thermal stability, 3 to 30% by volume of retained austenite is causedto be contained as the second phase. When the retained austenite contentis lower than 3% by volume, the elongation characteristics may possiblybe impaired and, when it is in excess of 30%, the thermal stability maypossibly be impaired. The volume fraction of retained austenite to becontained as the second phase is preferably 5 to 25%.

The second phase other than ferrite may further comprise, in addition tothose mentioned above, trace amounts, not larger than 1% by volume, ofcarbides, nitrides and oxides. As such, there may be mentionedcarbonitrides of Ti, Nb, V and Mo.

(C) Re: Grain Growth Rate at High Temperatures

The temperature characteristics of the grain growth rate in steel sheetsin which the average ferrite crystal grain diameter is within a specificrange so as to satisfy the relations defined by the formulas (1) and (2)given hereinabove depend on the rate of ferrite grain growth attemperatures in the vicinity of 700° C. Therefore, it becomes necessarythat the increase rate X (μm/min) in average ferrite crystal diameter at700° C. at the depth of ¼ of the sheet thickness from the steel sheetsurface and the above-mentioned average crystal grain diameter D (μm)satisfy the relation defined by the following formula (3):

D·X≦0.1  formula (3)

Thus, when the product D·X (μm²/min) of the increase rate X (μm/min) inaverage ferrite crystal grain diameter and the average crystal graindiameter D (μm) is maintained at a level not higher than 0.1 μm²/min,the steel sheet becomes stable against the main thermal history duringwelding or hot-dip plating and, accordingly, good thermal stability canbe obtained. For better thermal stability, the product D·X is preferablymaintained at 0.07 μm²/min or below, more preferably at 0.05 μm²/min orbelow.

As shown in Examples 2 and 3 given later herein, the ferrite crystalgrain microstructure of a steel sheet for which the product D·X(μm²/min) of the increase rate X (μm/min) in average ferrite crystalgrain diameter and the average crystal grain diameter D (μm) is nothigher than 0.1 μm²/min shows almost no changes in grain diameter evenupon tens of seconds of treatment at 850° C. The ferrite crystal grainsize (diameter) of the steel sheet according to the present inventionincreases almost proportionally to the time at 700° C., unlike theordinary grain growth which is proportional to the square root of thetime. Therefore, the increase rate X (μm/min) in average ferrite crystalgrain diameter is determined by measuring changes in grain diameter at700° C. for about 1 hour and calculating the mean rate of change.

In order to further decrease the grain growth rate, the dislocationdensity in ferrite crystal grains is kept preferably at a level nothigher than 10⁹/cm², more preferably not higher than 10⁸/cm².

(D) Re: Zn Plating:

The fine-grained hot-rolled steel sheet having the above-mentionedmicrostructure and thermal stability can be provided with a coat of Zn,Zn—Al alloy, Al—Si alloy or Fe—Zn alloy, for instance, on the steelsheet surface using a hot-dip plating line.

As for the Zn—Al alloy plating bath composition, a Zn-(0.1 to 60%) Albath, a composite bath derived by further addition of Si and/or Mg, orthe like is used. As for the Al—Si alloy plating bath composition, anAl-(7 to 13%) Si bath, for instance, is used. In the plating bath, theremay be contained at most 0.1% of Fe, V, Mn, Ti, Nb, Ca, Cr, Ni, W, Cu,Pb, Sn, Cd and/or Sb, without causing any trouble. The coat film on thesteel sheet surface after cooling following plating generally has acomposition somewhat higher in Fe concentration as compared with theplating bath composition since mutual diffusion of elements occursbetween the steel material and the molten metal during dipping andcooling. The technique of hot-dip alloying zinc plating positivelyutilizes this mutual diffusion, and the Fe concentration in the coatfilm amounts to 7 to 15%. The amount of the coating metal is notparticularly restricted but preferably is 30 to 200 g/m² per sidesurface; in the case of hot-dip alloying zinc plating, in order to avoida possibility of powdering, the amount is preferably 25 to 60 g/m².

The method of plating using a hot-dip plating line is as describedbelow.

The hot-rolled steel sheet after attaining a fine-grained microstructureis passed through a pickling step for scale removal from the surfacelayer and then enters a continuous hot-dip zinc plating line. Afteralkali degreasing and washing with water, in that order from the entryside, the steel sheet is preheated and then heated at a temperature of550 to 900° C. in a hydrogen-containing atmosphere for the reduction ofFe oxides on the steel sheet surface to prepare a surface suited for thesubsequent plating treatment. At temperatures below 550° C., theprogress of reduction is insufficient and, on the other hand, heating toa temperature exceeding 900° C. results in coarsening of the ferriticmicrostructure. When a ferrite+pearlite microstructure or aferrite+cementite microstructure is desired after plating, a temperaturerange of from 550° C. to about 730° C. is preferably employed. On theother hand, when bainite, martensite or retained austenite, forinstance, is to be formed, the temperature is preferably raised to atemperature range from the A₁ point to 900° C. in which the two phases,namely ferrite and austenite, can coexist. The hydrogen content in theatmosphere is preferably 5 to 40%. When the hydrogen content is lowerthan 5%, the reduction proceeds only to an insufficient extent. Atlevels exceeding 40%, the atmospheric gas cost becomes unduly high. Thecomponent other than hydrogen may be any gas that will not suppress thereduction. From the cost viewpoint, nitrogen is preferred. The period ofsoaking is only required to be such that the reduction can proceedsufficiently; hence it is not particularly specified but generally it isnot shorter than 10 seconds. The upper limit is 5 minutes at thelongest, more preferably 2 minutes at the longest, so that ferritecoarsening may be avoided. After passing the heating and soaking zonefor this reduction, the steel sheet is cooled to a temperature in thevicinity of the plating bath temperature, dipped in the plating bathand, after adjustment to a predetermined coating weight, cooled to roomtemperature. In the case of alloying hot-dip zinc plating, the steelsheet after hot-dip zinc plating in the above manner is reheated to 470to 600° C. so that the substrate iron and the coat film may react witheach other to form an Fe—Zn alloy layer on the steel sheet surface.

In this way, the steel sheet subjected to hot-dip plating is not onlyheated in the plating bath but also heat-treated at high temperatures inthe surface oxide layer reduction step prior to dipping in the platingbath and in the alloying step after dipping in the plating bath. Since,however, the steel sheet according to the present invention has athermally stable ferritic microstructure, it retains the fine-grainedmicrostructure and shows good mechanical characteristics even afterpassing those steps. Furthermore, the ferrite grains on the surface arefine, so that the alloying reaction rate increases and, therefore, theplated sheet can be produced efficiently and advantageously.

When it is to be plated, the steel preferably has a composition suchthat the C content is 0.001 to 0.15%, the Si content is 0.005 to 1.5%and/or P content is 0.005 to 1.0%.

(E) Re: Weldability:

The steel sheets having a fine-grained microstructure produced by theprior art low-temperature rolling are inferior in thermal stability andallow the HAZ to become softened and, therefore, the characteristics ofthe weld become deteriorated. On the contrary, the thermal stability ofthe steel sheet according to the invention is still good, even afterjoining by welding, in the form of the steel sheet as such orsurface-coated in the above manner to itself or to another member, thusimproving the formability of the weld after welding by the laser, spotor arc welding method, for instance.

As regards the chemical composition of the steel sheet to be subjectedto fusion welding, typically arc plasma welding or laser welding, it ispreferable to prescribe that the carbon equivalent Ceq(I) defined by theformula (7) given below should be 0.06 to 0.6%:

Ceq(I)=C+Mn/6+Si/24+Cr/5+Mo/4+Ni/40+V/14  formula (7)

Here, the symbols of elements in the above formula represent thecontents (in % by mass) of the respective elements in the steel.

Ceq(I) is an indicator of the maximum hardness of the weld and, byprescribing that Ceq(I) should be 0.06 to 0.6%, the formability of themember after welding can be secured. When Ceq(I) is smaller than 0.06%,the hardenability is poor and, therefore, the hardness of the weldedmetal portion becomes softer than the hardness of the base metalstrengthened by a thermally stable fine-grained microstructure and theworkability of the weld becomes decreased. And, when it is in excess of0.6%, the hardening, upon quench hardening, in the welded metal part andthe HAZ site having thermal stability is significant as compared withthe base metal hardness and, therefore, the formability of the welddecreases. It is preferable to prescribe that Ceq(I) should be 0.10 to0.5%. Further, it is preferable to prescribe that the content of C,which causes hardening and embrittlement of the weld, should be nothigher than 0.17% by mass.

In spot welding in which resistance welding is realized byelectricity-based heat generation in the base metal members, it ispreferable to prescribe, from the viewpoint of hardness distribution inthe weld and prevention of embrittlement in order to secure the jointstrength, that, regarding the chemical composition, the C content shouldbe not higher than 0.17% by mass and the carbon equivalent Ceq(II)defined by the formula (8) given below should be 0.03 to 0.20% and,further, in order to obtain a nugget diameter for securing the jointstrength within a broad range of conditions, that the indicator Rsp ofthe base metal resistance defined by the formula (9) given below shouldbe not higher than 45:

Ceq(II)=C+Mn/100+Si/90+Cr/100  formula (8)

Rsp=13.5×(Si+Al+0.4Mn+0.4Cr)+12.2  formula (9)

Here, the symbols of elements in the above formulas represent thecontents (in % by mass) of the respective elements in the steel.

In a thermal cycle involving rapid cooling, as in spot welding, theinfluence of the C content on hardening and embrittlement is great, sothat the C content is preferably not higher than 0.17%, more preferablynot higher than 0.15%.

Ceq(II) is an indicator of the maximum hardness in the weld in such arapid cooling-involving thermal cycle as in spot welding and, byprescribing that Ceq(II) should be 0.03 to 0.20%, it becomes possible toobtain a weldable part excellent in resistance weldability. When Ceq(II)is lower than 0.03%, the hardenability is poor, the maximum hardnessobtainable in the spot-welded portion is low as compared with the basemetal hardness and, therefore, the so-called button breakage can beobtained in joint strength evaluation testing but the maximum breakingload attainable is at a low level. When Ceq(II) is in excess of 0.20%,the hardening and embrittlement due to quench hardening in the weldmetal portion and the HAZ site having thermal stability are remarkableas compared with the base metal hardness and, in strength evaluationtesting, cracking occurs in the molten metal site (within nugget) and itbecomes difficult to attain the so-called button breakage. Preferably,it is prescribed that Ceq(II) should be 0.06 to 0.17%.

The indicator Rsp of base metal resistance is indicative of how broad isthe range of welding conditions under which a nugget (fusion weld site)diameter for securing the joint strength can be obtained and, in orderto obtain a weldable site excellent in resistance weldability, it ispreferably not higher than 45, more preferably not higher than 40.

In order to obtain a nugget diameter for securing the joint strengthwithin a broad range of welding conditions, the current density andresistance heating become important factors. Here, the current densityis determined by the sectional area of the route of current flow duringwelding and, in the case of the steel according to the present inventionwhich is excellent in thermal stability, the softening due to graingrowth will not occur, so that the initial route of current flow isprevented from broadening and a sufficient nugget diameter is readilyformed. On the other hand, the resistance heating is greatly influencedby the electric resistance of the base metal; when the base metalresistance is high, excessive resistance heating will occur and, when itis in excess of the optimum condition range, expulsion and surface flashwill readily occur. Generally, the optimum condition range in spotwelding is expressed in terms of the range from the welding currentcausing formation of a nugget diameter of 4×√{square root over ( )}tbeing the jointing material sheet thickness) to the expulsion current orin terms of the range from the minimum current for button breakage beingmanifested to the expulsion current.

(F) Re: Rolling:

The rolling is carried out from a temperature exceeding 1000° C. in theaustenite temperature range using a reversing mill or tandem mill unit.From the industrial productivity viewpoint, a tandem mill unit ispreferably used in at least several final stages.

A slab obtained by continuous casting or casting and slabbing or a steelsheet obtained by strip casting, for instance, if necessary oncesubjected to hot or cold working, is used and rolled after reheating toa temperature exceeding 1000° C., if it is a cold one. When the rollingstart temperature is not higher than 1000° C., the rolling load becomesexcessive and it becomes difficult to attain a sufficient reduction and,in addition, it becomes difficult to finish the rolling at a temperaturenot lower than the Ar₃ point with a sufficient reduction, with theresult that the desired mechanical characteristics or thermal stabilitycannot be obtained. Preferably, the rolling is started from atemperature not lower than 1025° C., more preferably not lower than1050° C. In order to prevent austenite grains from becoming coarse andto suppress the equipment cost and heating fuel cost, the upper limit isset at a level not higher than 1350° C., preferably not higher than1250° C. In the case of steel species for which there is no need tosufficiently dissolve such a precipitate as TiC or NbC in austenite, itis preferable to reheat the material to a relatively low temperature(1050 to 1150° C.) within the range mentioned above. This is becauseinitial austenite crystal grains are rendered fine and final ferritecrystal grains also become capable of being readily rendered fine.

The rolling finish temperature is selected within the range of not lowerthan the Ar₃ point and not lower than 780° C. so that austenite may betransformed into ferrite after rolling. When the finish temperature islower than the Ar₃ point, ferrite is formed during rolling. When thefinish temperature is lower than 780° C., the rolling load increases andit becomes difficult to apply a sufficient reduction and, in addition,ferrite transformation may occur in the sheet surface layer duringrolling. Preferably, the rolling is finished at a temperature not lowerthan the Ar₃ point and not lower than 800° C.

The rolling finish temperature is preferably as low as possible withinthe range not lower than the Ar₃ point and not lower than 780° C. Thisis because the effect of accumulating work strains introduced intoaustenite by rolling is increased and the reduction in size of crystalgrains is promoted thereby. The Ar₃ point of the steel species to beused in the practice of the invention is generally in the range of 780to 900° C.

The total reduction, as expressed in terms of sheet thickness reduction,in order to promote the size reduction of ferrite is not lower than 90%,preferably not lower than 92%, more preferably not lower than 94%. Thesheet thickness reduction within the temperature range from the rollingfinish temperature to the “rolling finish temperature+100° C.” ispreferably not lower than 40%. More preferably, the sheet thicknessreduction within the temperature range from the rolling finishtemperature to the “rolling finish temperature+80° C.” is not lower than60%. The rolling is carried out in a plurality of continuous passes. Thereduction per pass is preferably 15 to 60%. A higher reduction per passis preferred from the viewpoint of accumulating strains imposed onaustenite and rendering the crystal grain diameter of ferrite formed bytransformation but it re-quires an increase in rolling load, which inturn requires a large-sized rolling plant and in addition makes itdifficult to control the sheet morphology. In accordance with the methodaccording to the present invention, fine ferrite crystal grains can beobtained even in the rolling including a plurality of passes with areduction per pass of not higher than 40%. Therefore, when it is desiredthat the sheet morphology control be facilitated, it is preferable tocarry out the final two passes at a reduction per pass of not higherthan 40%.

(G) Re: Cooling after Rolling:

After finishing the rolling, the sheet is cooled to a temperature nothigher than 720° C. within 0.4 second from the moment of finishing therolling so that work strains introduced into ferrite may not be relievedbut may serve as driving forces for the transformation of austenite toferrite to form a microstructure comprising fine ferrite crystal grainsPreferably, the sheet is cooled to a temperature not higher than 720° C.within 0.2 second from the moment of finishing the rolling As for thecooling, water cooling is desirably employed, and the rate of cooling ispreferably not lower than 400° C./second as expressed in terms ofaverage cooling rate for the period during which forced cooling iscarried out, excluding the period of air cooling.

The reason why the time until cooling to a temperature not higher than720° C. is prescribed here is that when the cooling is discontinued orslowed down at a temperature exceeding 720° C., the strains introducedby working are relieved before fine ferrite grains are formed, or thestrains change in their state of occurrence and no more effectivelyserve as nuclei for ferrite, so that ferrite crystal grains becomeremarkably coarse.

When the temperature arrives at a level not higher than 720° C., thesheet enters the transformation temperature range in which the ferritetransformation becomes active. The ferrite transformation temperaturerange in which the above-mentioned ferrite microstructure is obtainedranges from that temperature to 600° C. Therefore, after arrival at 720°C. or below, the cooling is interrupted or the rate of cooling islowered, and the sheet is maintained in this temperature range for atleast 2 seconds, whereby the formation of the above-mentioned thermallystable ferrite crystal grain microstructure can be secured. When theholding time is shorter, the formation of the above-mentioned thermallystable ferrite crystal grain microstructure may possibly be suppressed.More preferably, the sheet is allowed to remain in the temperature rangeof 620 to 700° C. for at least 3 seconds.

In order to produce a double phase microstructure steel species the mainphase of which is a fine ferrite crystal grain microstructure and whichcontains at least 5% by volume of martensite dispersed therein, thetemperature is preferably lowered to 350° C. or below after theabove-mentioned cooling and residence. More preferably, the sheet iscooled to a temperature of 250° C. or below at a cooling rate of notlower than 40° C./sec. If the cooling to a temperature of 350° C. orbelow is carried out at a cooling rate not higher than 20° C./sec,bainite will be formed readily and the martensite formation may possiblybe suppressed.

On the other hand, in order to produce a double phase microstructuresteel species the main phase of which is a fine ferrite crystal grainmicrostructure and which contains 3 to 30% by volume of retainedaustenite dispersed therein, it is preferred that, after theabove-mentioned cooling, the sheet be cooled to 350 to 500° C. at acooling rate of not lower than 20° C./sec and thereafter cooled slowlyat a cooling rate of not higher than 60° C./hour. It is more preferablethat the rate of cooling to 400 to 500° C. be not lower than 50° C./sec.

(H) Re: Cooling Equipment:

In carrying out the present invention, the equipment for carrying outthe cooling mentioned above is not restricted. Industrially, the use ofa water-spraying apparatus with a high water volume density isappropriate. For example, cooling can be effected by disposing waterspray headers between the rolled sheet conveying rollers and sprayinghigh pressure water from above and below the sheet at a sufficient watervolume density.

(I) Re: Cold Rolling and Annealing:

In order to efficiently produce a thin steel sheet having a fine-grainedmicrostructure, the sheet after hot rolling is pickled, then furthercold-rolled and, thereafter, annealed. The reduction in cold rolling isnot lower than 40% in order to promote the recrystallization of ferriteduring annealing but not higher than 90% considering the fact that therolling becomes difficult to perform. The rolling equipment is notrestricted but a tandem mill or reversing mill unit is preferably used.

In order to attain a fine-grained ferritic microstructure byrecrystallization of worked ferrite after cold rolling, the sheet issubjected to heat treatment. The temperature is not lower than thetemperature at which ferrite recrystallization occurs and not higherthan 900° C. so that crystal grains are suppressed from becoming coarse.Preferably, it is not lower than the Ac₁ point and not higher than theAc₃ point. At below the Ac₁ point, a long period of time is required forferrite recrystallization and, at above the Ac₃ point, themicrostructure becomes an austenite single phase, so that themicrostructure tends to become coarse. The annealing time is not shorterthan the time required for ferrite recrystallization; the upper limit isnot restricted. Ordinary continuous annealing equipment or batchwiseannealing equipment, for instance, can be used; for efficientproduction, however, it is preferable to use continuous annealingequipment and carry out the annealing in a short period of time. Whenhot-dip plating is carried out using continuous hot-dip platingequipment, the plating equipment generally includes a pre-annealingstep, so that it is not necessary to carry out annealing after coldrolling but the cold-rolled material can be directly submitted to theplating equipment.

The following examples illustrate the present invention in furtherdetail.

Example 1

Steel species Al to All having the respective chemical compositionsshown in Table 1 were prepared by melting and subjected to hot forgingto give 30 mm-thick plates. Thereafter, they were reheated to 1050° C.or above and then rolled on a small-sized test tandem mill to give 2mm-thick sheets.

TABLE 1 Steel species C Si Mn P S Al N Other components Ar₃ 2.7 +5000/(5 + 350 · C + 40 · Mn)² A1 0.105 0.05 1.03 0.060 0.0035 0.0300.0043 824 3.43 A2 0.075 0.03 0.44 0.021 0.0025 0.030 0.0030 865 4.80 A30.030 0.03 0.22 0.092 0.0031 0.030 0.0027 890 11.17 A4 0.048 0.03 0.300.063 0.0023 0.015 0.0035 Ca: 0.0021 881 7.08 A5 0.155 0.03 0.62 0.0220.0041 0.030 0.0024 Nb: 0.010 824 3.41 A6 0.080 0.06 0.70 0.018 0.00560.034 0.0030 Ni: 0.29, Cr: 0.31 836 4.04 A7 0.061 0.54 1.53 0.015 0.00250.025 0.0041 Mo: 0.21 821 3.35 A8 0.080 0.03 1.00 0.016 0.0030 0.0350.0035 Nb: 0.003, Ti: 0.013, Ca: 0.0021 836 3.64 A9 0.050 0.01 1.800.018 0.0038 0.037 0.0030 Ti: 0.022 810 3.26  A10 0.045 0.42 0.70 0.0250.0042 0.043 0.0030 V: 0.1 882 4.80  A11 0.023 0.02 0.18 0.011 0.00320.032 0.0028 893 14.89

The rolling finish temperatures and cooling conditions were as shown inTable 2. In each rolling process, the rolling finish temperature washigher than the Ar₃ point of the corresponding steel species andmulti-pass rolling was carried out in at least 3 passes within thetemperature range of the finish temperature to the [finishtemperature+100° C.]. In the final two passes, low-reduction rolling wasconducted at a reduction per pass of 35%, except for Test No. 3. In thefinal two passes in Test No. 3, high-reduction rolling was carried outat 50 to 60%. After rolling finished, water cooling was conducted to apredetermined temperature within the temperature range of 500 to 720°C., as shown in Table 2. In some Test Numbers, the water cooling wasfollowed by air cooling and thus a holding time was provided in order tomaintain the sheet in the range of 720 to 600° C. In Table 2, there areshown the holding time in the temperature range of 700 to 620° C. inaddition to the holding time in the temperature range of 720 to 600° C.Thereafter, water cooling to room temperature was carried out at a rateof about 100° C./sec or, after water cooling to a predeterminedtemperature within the temperature range of 600 to 400° C., cooling wascarried out in a furnace, to give steel sheets varying in second phasemicrostructure.

TABLE 2 Holding time (s) in Temperature (° C.) at Finishing Cooling time(s) from Cooling Holding time (s) the temperature range of which 100°C./s rate Test Steel temperature completion of finishing rate in thetemperature 700 to 620° C. within the water cooling was stopped No.species (° C.) rolling to 720° C. (° C./s) range of 720 to 600° C.left-mentioned range at 600° C. or below 1 A1 835 0.14 795 0.8 0.6 500 2A1 839 0.15 817 3.4 3.1 RT 3 A1 765 0.05 1021 4.1 3.6 RT 4 A1 842 0.111093 11.1  10.0 500 5 A1 861 1.52 1080 3.2 3.0 500 6 A2 865 0.15 976 5.14.7 RT 7 A2 871 0.16 949 7.5 7.0 RT 8 A2 875 0.14 1123 6.5 5.9 500 9 A2881 0.15 1096 11.4  10.4 500 10 A3 892 0.13 1620 3.2 2.8 RT 11 A3 8940.19 1400 3.8 3.4 RT 12 A3 895 0.14 1500 5.8 5.3 RT 13 A3 901 0.17 124010.8  9.8 RT 14 A4 882 0.16 1019 5.3 4.8 RT 15 A4 878 0.17 939 5.7 5.2RT 16 A4 883 0.14 1163 3.4 3.2 500 17 A5 842 0.13 911 3.5 3.2 500 18 A5825 0.15 708 7.8 7.1 500 19 A5 850 0.16 920 0.2 0.2 RT 20 A6 845 0.16795 6.4 5.8 500 21 A7 833 0.11 1064 5.3 4.9 RT 22 A8 845 0.12 1050 5.34.9 RT 23 A9 820 0.10 977 7.3 6.7 RT 24  A10 892 0.19 919 4.3 4.0 RT 25 A11 895 0.14 1650 6.9 6.2 RT Note: Each underline indicates that therelevant value is out of scope of the producing conditions of thepresent invention. RT: Room temperature

The microstructure of each of the thus-obtained hot-rolled steel sheetswas observed in cross section in the direction of steel sheet thicknessusing a scanning electron microscope.

The ferrite crystal grain diameter and the ferrite grain diameterdistribution at the depth of ¼ of the sheet thickness from the sheetsurface were determined by carrying out crystal orientation analysis bythe EBSP (electron back scattering pattern) method. Observing themicrostructure at the depth of ¼ of the sheet thickness from the sheetsurface as etched with nital or picric acid using a scanning electronmicroscope determined the volume fraction of each phase. Themicrostructure of the second phase other than the ferrite phase in thesteel sheets produced in this example comprised pearlite, bainite andintragranular spherical cementite or grain boundary cementite.

For each of the steel sheets according to the present invention, thecrystal grain diameter at the depth of 100 μm from the steel sheetsurface and the crystal grain diameter in the center were measured inthe same manner as mentioned above. As a result, in all the steel sheetsaccording to the present invention, the crystal grain diameter at thedepth of 100 μm was not greater than 60% of the grain diameter in thecenter of the sheet thickness, and the grain diameter at the depth of ¼of the sheet thickness was not greater than 85% of the grain diameter inthe center of the sheet thickness.

As for the mechanical properties, tensile characteristics were testedusing JIS No. 5 tensile test specimens, and the tensile strength TS(MPa), yield ratio YR and total elongation El (%) were evaluated.

As for the thermal stability, each specimen was immersed in a salt bathat 700° C. for 10, 30 or 60 minutes and then rapidly cooled and thegrain diameter was measured by the same method as mentioned above, andthe increase rate X (μm/min) in average crystal grain diameter wascalculated by dividing the difference between the grain diameter d₀ (μm)before annealing and the grain diameter d₁ (μm) after annealing by theannealing time (min).

In Table 3, there are shown the microstructure and properties of each ofthe thus-obtained hot-rolled steel sheets and the tensile test results.Here, in Test No. 1, the holding time in the temperature range of 720 to600° C. was as short as 0.8 second, so that the ferrite volume fractionwas as small as 14.8% and, further, the grain growth rate upon annealingat 700° C. was high; thus, the sheet was inferior in thermal stability.In Test No. 3 in which high reduction rolling at low temperatures wasemployed, the grain diameter was excessively small, namely 1.13 μm, andthe sheet was inferior in thermal stability and strength-elongationbalance. In Test No. 5, the cooling time from completion of finishingrolling to 720° C. was 1.52 seconds and, therefore, the average ferritecrystal grain diameter became 4.52 μm, indicating grain coarsening, andthe microstructure became a duplex grain microstructure; hence the sheetwas inferior in thermal stability. In Test No. 19, the holding time inthe temperature range of 720 to 600° C. was very short, namely 0.2second, so that the microstructure became such that the amount ofbainite was in excess of 95% and the ferrite volume fraction was as lowas below 5%. Contrary to these comparatives, the steel sheets obtainedaccording to the present invention in which the cooling conditions werewithin the range specified herein were superior in both thermalstability and mechanical properties.

TABLE 3 Increase rate X Average (μm/ ferrite min) D/3 ≦ Dis- crystal ingrain d ≦ 3D location grain diam- Area density Ferrite Second phasevolume fraction Mechanical properties diam- eter D · X per- (cm⁻²)volume Per- Mar- Granular TS × El Test eter at (μm²/ centage in ferritefraction lite Bainite tensite cementite TS YR El (MPa · No. D (μm) 700°C. min) (%) grains (%) (%) (%) (%) (%) (MPa) (—) (%) %) 1 2.23 0.2510.560 85 2.0 × 10⁹ 14.8 — 85.2 — — 578 0.96 26.5 15317 Comparative 22.38 0.010 0.024 92 8.5 × 10⁷ 88.5 1.0 8.50 2.0 — 563 0.94 30.7 17289Invention 3 1.13 0.110 0.124 75 6.5 × 10⁹ 94.4 — 5.0 — 0.6 654 0.98 15.310006 Comparative 4 2.62 0.009 0.024 91 9.5 × 10⁷ 89.4 10.4 — — 0.2 5050.91 33.0 16677 Invention 5 4.52 0.221 0.999 65 4.6 × 10⁷ 87.5 12.5 — —— 445 0.84 34.5 15330 Comparative 6 4.63 0.004 0.019 85 6.2 × 10⁷ 90.51.0 7.5 1.0 — 475 0.86 32.1 15248 Invention 7 4.17 0.005 0.020 91 7.6 ×10⁷ 92.2 4.2 2.6 1.0 — 485 0.86 32.5 15763 8 5.06 0.004 0.019 85 8.2 ×10⁷ 92.7 3.2 4.0 — 0.1 414 0.83 37.0 15301 9 4.76 0.004 0.019 86 4.1 ×10⁷ 93.4 4.2 2.3 — 0.1 407 0.83 39.8 16207 10 4.25 0.005 0.021 93.5 4.0× 10⁷ 95.6 1.0 3.4 — — 434 0.90 25.8 15057 11 5.32 0.006 0.032 96 5.8 ×10⁷ 97.0 1.5 1.5 — — 422 0.92 36.5 15400 12 4.50 0.005 0.023 96.5 5.6 ×10⁷ 96.7 2.1 1.2 — — 410 0.92 37.2 15250 13 4.93 0.006 0.030 98 4.1 ×10⁷ 96.4 2.2 0.9 0.5 — 395 0.91 41.3 16313 14 5.06 0.004 0.019 97 3.6 ×10⁷ 94.5 2.0 3.0 0.5 — 429 0.89 35.1 15049 15 4.76 0.004 0.019 95 6.6 ×10⁷ 94.3 1.5 3.2 1.0 — 433 0.91 36.2 15680 16 5.25 0.003 0.018 95 3.1 ×10⁷ 95.0 3.2 1.7 — 0.1 375 0.85 40.3 15092 17 2.95 0.008 0.023 96 7.8 ×10⁷ 80.6 17.2 2.0 — 0.2 600 0.88 25.6 15360 18 2.54 0.009 0.024 95 5.6 ×10⁷ 78.5 20.3 1.0 — 0.2 693 0.79 25.8 17879 19 *1 ≦5    — >95 — 0.2 9600.9  6.4 6143 Comparative 20 2.52 0.009 0.024 96 9.6 × 10⁷ 91.0 6.9 2.0— 0.1 488 0.92 32.1 15671 Invention 21 1.65 0.017 0.028 88 7.9 × 10⁷92.6 2.0 5.4 — — 588 0.97 26.1 15338 22 2.75 0.008 0.023 98 5.2 × 10⁷91.7 — 6.3 2.0 — 535 0.90 32.3 17293 23 2.13 0.012 0.025 96 6.6 × 10⁷93.3 — 2.2 4.5 — 593 0.96 29.6 17544 24 3.62 0.006 0.021 95 4.9 × 10⁷94.8 — 5.2 — — 549 0.88 29.5 16194 25 6.50 0.012 0.078 86 3.1 × 10⁷ 98.0— 2.0 — — 423 0.76 37.6 15905 *1: The ferrite volume fraction was toosmall to be determined.

Example 2

Semi-finished products (size: 80 mm in width×100 mm in length×35 mm inthickness) respectively made of steel species 1 to 5 having therespective chemical compositions shown in Table 4 were hot-rolled at atemperature not lower than the Ar₃ point under the conditions shown inTable 5, followed by water cooling, to give hot-rolled steel sheets witha sheet thickness of 1.2 mm.

TABLE 4 Steel species C Si Mn P S Al N Ar₃ 2.7 + 5000/(5 + 350 · C + 40· Mn)² 1 0.15 0.01 0.75 0.010 0.002 0.035 0.003 827 3.35 2 0.10 0.021.02 0.011 0.002 0.035 0.004 834 3.47 3 0.19 0.03 0.99 0.012 0.001 0.0300.002 803 3.11 4 0.07 0.05 0.75 0.012 0.003 0.034 0.003 855 4.11 5 0.100.02 1.40 0.009 0.002 0.031 0.004 819 3.24

TABLE 5 Holding time (s) in the Cooling time (s) Holding time (s)temperature range of Finishing from completion of in the temperature 700to 620° C. within Time (s) to cooling Steel temperature finishingrolling to Cooling rate range the left-mentioned to 250° C. after TestNo. species (° C.) 720° C. (° C./s) of 720 to 600° C. range holding A 1830 0.17 987 3.3 1.9 8.8 Invention B C 756 0.17 945 1.6 0.56 12.3Comparative D 2 843 0.20 1201 3.0 1.6 11.5 Invention E 3 814 0.17 11113.4 2.3 9.1 F 771 0.17 826 1.5 0.67 13.3 Comparative G 4 861 0.20 10234.0 2.7 10.9 Invention H 5 825 0.20 1132 3.1 1.8 12.6

The hot-rolled steel sheets obtained were measured for average ferritecrystal grain diameter and ferrite grain diameter distribution and fordislocation density and evaluated for thermal stability. Thus, theferrite crystal grain diameter, ferrite grain diameter distribution anddislocation density were measured, and the thermal stability evaluationwas performed in the same manner as mentioned above. The dislocationdensity ρ (cm⁻²) was determined by measuring the number N of points ofintersection and contact between an arbitrary line segment having alength L (cm) and dislocation lines in the bright field image under atransmission electron microscope and making a calculation according tothe formula (10) given below:

ρ=2N/Lt  formula (10)

wherein t is the sheet thickness (cm).

As for the thermal stability of ferrite crystal grains, each specimenwas immersed in a salt bath at 700° C. for 10, 30 or 60 minutes and thenrapidly cooled and the grain diameter was measured by the same method asmentioned above, and the increase rate X (μm/min) in average crystalgrain diameter was calculated by dividing the difference between thegrain diameter d₀ (μm) before annealing and the grain diameter d₁ (μm)after annealing by the annealing time (min).

For each of the steel sheets according to the present invention, thecrystal grain diameter at the depth of 100 μm from the steel sheetsurface and the crystal grain diameter in the center were measured inthe same manner as mentioned above. As a result, in all the steel sheetsaccording to the present invention, the crystal grain diameter at thedepth of 100 μm was not greater than 60% of the grain diameter in thecenter of the sheet thickness, and the grain diameter at the depth of ¼of the sheet thickness was not greater than 85% of the grain diameter inthe center of the sheet thickness.

The thus-obtained hot-rolled steel sheets were subjected to reheatingtreatment within the range of 730 to 830° C. to reveal the influencesexerted by the reheating treatment on the mechanical properties of thesteel sheets, and the average ferrite crystal grain diametermeasurements were again carried out. As for the mechanical properties,the tensile characteristics were tested using JIS No. 5 tensile testspecimens, and the tensile strength TS (MPa), yield ratio YR and totalelongation El (%) were evaluated.

In Table 6, there are shown, for each hot-rolled steel sheet obtained inthe above manner, the microstructure, the properties and the tensiletest results of as well as the results of the second average ferritecrystal grain diameter measurement after the reheating treatment withinthe range of 730 to 830° C. Here, in Test Nos. C and F, the steel sheetsafter hot rolling were inferior in mechanical characteristics and inthermal stability as well and, after the reheating treatment, theferrite crystal grain diameters were in excess of 8 μm and furtherdeteriorations in mechanical characteristics could be confirmed. Incontrast to these comparatives, the steel sheets excellent in thermalstability in the examples of the practice of the present inventionshowed excellent mechanical characteristics and showed almost no changesin grain diameter even after tens of seconds of reheating treatment at730 to 830° C. Thus, it could be confirmed that the steel sheetsaccording to the present invention were found fine grain-strengthenedeven after the reheating treatment.

TABLE 6 Average Increase ferrite rate X Dislocation crystal (μm/min)density D/3 ≦ d ≦ 3D Ferrite Second phase volume fraction grain in grain(cm⁻²) Area volume Granular Test Steel diameter diameter D · X inferrite percentage fraction Perlite Bainite Martensite cementite No.species D (μm) at 700° C. (μm²/min) grains (%) (%) (%) (%) (%) (%) A 11.9 0.015 0.029 9.5 × 10⁷ 93 92.3 6.5 — — 1.2 B C 1.5 0.12 0.18   1.1 ×10¹⁰ 72 90.7 2.4 6.9 — — D 2 2.0 0.021 0.042 9.1 × 10⁷ 96 95.7 1.7 — —2.6 E 3 1.6 0.012 0.019 1.2 × 10⁸ 95 90.5 8.9 — — 0.6 F 1.3 0.13 0.17  4.5 × 10¹⁰ 68 92.4 — — — 7.6 G 4 2.3 0.024 0.055 6.3 × 10⁷ 92 97.3 — —— 2.7 H 5 1.8 0.016 0.029 8.8 × 10⁷ 95 95.1 — 2.4 — 2.5 Reheating andinfluences thereof Average ferrite crystal Mechanical propertiesReheating grain Test TS YR El TS × El temperature Holding diameter No.(MPa) (—) (%) (MPa · %) (° C.) time (s) D (μm) A 575 0.97 26.5 15238 78030 1.9 Invention B 830 10 2.0 C 599 0.99 13.7 8206 830 10 8.2Comparative D 545 0.96 28.1 15315 750 60 2.1 Invention E 693 0.95 23.416216 830 10 1.6 F 705 0.99 11.0 7755 780 60 11.6 Comparative G 575 0.9627.5 15813 750 30 2.3 Invention H 556 0.96 28.3 15735 750 30 1.8

Example 3

Steel species AA to AZ having the respective chemical compositions shownin Table 7 were prepared by melting, followed by hot forging to reducethe thickness to 30 mm. Thereafter, each semi-finished product wasreheating ed to a temperature of 1100 to 1200° C. and then rolled in 5passes at a temperature higher than the Ar₃ point to give a 2 mm-thickfinished sheet. In the last two passes, light reduction rolling wascarried out at a reduction not higher than 35%/pass. After rolling, eachsheet was cooled under the conditions shown in Table 8. Themicrostructure of each steel material obtained was observed in a crosssection of the steel sheet thickness using a scanning electronmicroscope (SEM).

TABLE 7 Steel species C Si Mn P S Al Cr Mo Nb Ti 2.7 + 5000/(5 + 350 ·C + 40 · Mn)² Ar₃ (° C.) AA 0.05 0.03 2.49 0.011 0.003 0.034 — — — —3.04 775 AB 0.09 1.08 1.80 0.018 0.001 0.030 — — 0.010 — 3.13 820 AC0.06 0.80 1.58 0.015 0.002 0.032 0.48 — — — 3.33 830 AD 0.06 0.49 1.980.012 0.003 0.030 0.26 — 0.009 — 3.15 800 AE 0.07 0.51 1.56 0.014 0.0010.033 0.62 — — 0.013 3.29 815 AF 0.08 0.58 2.10 0.013 0.003 0.031 — — —— 3.07 800 AG 0.05 0.53 1.56 0.016 0.002 0.036 — 0.24 — — 3.39 835 AY0.28 0.04 2.45 0.014 0.002 0.034 — — — — 2.82 730 AZ 0.06 0.03 3.360.012 0.003 0.037 — — — — 2.89 740

TABLE 8 Holding time (s) in Holding time (s) the temperature Time (s) toFinishing Cooling time (s) from Cooling in the temperature range of 700to 620° C. cooling to Test Steel temperature completion of finishingrate range within the 250° C. after No. species (° C.) rolling to 720°C. (° C./s) of 720 to 600° C. left-mentioned range holding A1 AA 8000.065 1231 4.3 4.1 73 A2 AB 845 0.10  1250 6.6 6.4 78 A3 AC 850 0.11 1182 6.5 6.2 81 A4 AD 820 0.085 1176 4.2 3.9 72 A5 AE 835 0.094 1223 6.05.8 78 A6 AF 820 0.085 1176 2.7 2.5 74 A7 AG 860 0.12  1167 7.4 7.1 68A8 AY 800 0.067 1194 8.0 7.8 56 A9 AZ 800 0.072 1111 9.8 9.6 57 A10 AA800 0.58   138 3.8 3.5 76 A11 AD 800 0.17   471 1.8 1.6 78 Note: Eachunderline indicates that the relevant value is out of scope of theproducing conditions of the present invention.

The ferrite crystal grain diameter and grain diameter distribution weredetermined by carrying out crystal orientation analysis at the depth of¼ of the sheet thickness from the sheet surface using the EBSP (electronback scattering pattern) method. Observing the microstructure at thedepth of ¼ of the sheet thickness from the sheet surface as etched withnital or picric acid using a SEM determined the volume fraction of eachphase. The ferrite volume fraction and martensite volume fraction weremeasured at the depth of ¼ of the sheet thickness from the sheet surfaceby the so-called mesh method and were expressed in terms of arithmeticmeans. Further, JIS No. 5 test specimens were taken from each rolledmaterial and the mechanical characteristics were evaluated on a tensiletester at ordinary temperature.

For each of the steel sheets according to the present invention, thecrystal grain diameter at the depth of 100 μm from the steel sheetsurface and the crystal grain diameter in the center were measured inthe same manner as mentioned above. As a result, in all the steel sheetsaccording to the present invention, the crystal grain diameter at thedepth of 100 μm was not greater than 60% of the grain diameter in thecenter of the sheet thickness, and the grain diameter at the depth of ¼of the sheet thickness was not greater than 85% of the grain diameter inthe center of the sheet thickness.

As for the thermal stability, each specimen was immersed in a salt bathat 700° C. for 10, 30 or 60 minutes and then rapidly cooled, and theincrease rate X (μm/min) in average crystal grain diameter wascalculated by the same method as described above.

These results are shown in Table 9. Here, in Test No. A10, the ferritecrystal grain diameter was as coarse as 4.57 μm, so that the sheet wasinferior in mechanical characteristics and in thermal stability. In eachof Test Nos. A8, A9 and All, the product D·X of the increase rate X inaverage crystal grain diameter and the average crystal grain diameter Dwas in excess of 0.1 μm²/min and the ferrite volume fraction was small,so that the sheet was inferior in mechanical characteristics and inthermal stability. In contrast with these comparatives, the steel sheetsobtained in Test Nos. A1 to A7 in accordance with the present inventioneach had a fine martensite crystal grain diameter around 2.5 μm in spiteof light reduction rolling and comprised 50% by volume or more offerrite and 10% by volume or more of martensite. Such ferritemicrostructures are thermally stable and contain an appropriate amountof martensite and therefore can show high strength and good elongationcharacteristics.

TABLE 9 Average Increase ferrite rate X D/3 ≦ crystal (μm/min) d ≦ 3DMechanical properties grain in grain D · X Area Dislocation TS × El TestMicro- diameter diameter (μm²/ percentage density (cm⁻²) FerriteMartensite TS YR El (MPa · No. structure D (μm) at 700° C. min) (%) inferrite grains (%) (%) (MPa) (—) (%) %) A1 α + M 2.64 0.010 0.026 89 7.2× 10⁷ 65.7 34.3 813 0.62 20.5 16667 Invention A2 α + M 2.51 0.009 0.02388 5.4 × 10⁷ 78.2 21.8 938 0.61 20.6 19323 A3 α + M 2.76 0.010 0.028 935.7 × 10⁷ 88.6 11.4 789 0.63 24.1 19015 A4 α + M 2.83 0.008 0.022 90 6.1× 10⁷ 75.8 24.2 788 0.58 24.5 19306 A5 α + M 2.48 0.009 0.022 88 5.6 ×10⁷ 78.9 21.1 748 0.55 25.9 19373 A6 α + M 2.62 0.010 0.025 91 6.9 × 10⁷74.8 25.2 804 0.65 23.8 19135 A7 α + M 2.75 0.011 0.029 90 5.3 × 10⁷77.9 22.1 791 0.59 24.8 19617 A8 α + M + B 2.47 0.048 0.118 82 8.6 × 10⁷37.8 37.1 1027 0.81 10.2 10475 Comparative A9 α + M + B 2.39 0.054 0.12873 1.9 × 10⁹ 28.8 40.7 965 0.84 8.9 8589 A10 α + M + B 4.57 0.010 0.04687 8.8 × 10⁷ 41.2 29.5 786 0.82 10.5 8253 A11 α + M 2.98 0.037 0.109 822.1 × 10⁹ 46.2 53.8 615 0.78 22.7 13961 Note: The symbols in the columnof microstructure indicate the following meaning: α: ferrite, M:Martensite, B: Bainite.

Example 4

Steel species A1 to A10 having the respective chemical compositionsshown in Table 10 were prepared by melting, followed by hot forging toreduce the thickness to 35 mm. Thereafter, each semi-finished productwas reheated to a temperature of 1050 to 1250° C. and then rolled in 5passes at a temperature higher than the Ar₃ point to give a 1.5 mm-thickfinished sheet. After rolling, each sheet was cooled under theconditions shown in Table 11. The microstructure of each steel materialobtained was observed in a cross section of the steel sheet thicknessusing a scanning electron microscope (SEM) method.

TABLE 10 Steel Nb + 2.7 + 5000/ species C Si Mn Al P S Nb Ti Ti (5 + 350· C + 40 · Mn)² Ar₃ (° C.) A1 0.19 0.46 2.45 0.84 0.008 0.002 — — — 2.87825 A2 0.14 0.46 1.98 0.83 0.012 0.001 — — — 2.98 868 A3 0.20 0.49 2.020.87 0.010 0.003 0.008 — 0.008 2.91 844 A4 0.09 0.57 1.98 0.76 0.0080.006 — — — 3.07 893 A5 0.11 1.75 1.22 0.36 0.025 0.008 0.013 0.0150.028 3.29 879 A6 0.11 1.32 1.26 0.35 0.015 0.001 — 0.019 0.019 3.27 891A7 0.12 0.75 1.83 0.57 0.018 0.003 — — — 3.05 854 A8 0.14 1.48 1.66 0.020.008 0.001 0.013 — 0.013 3.04 817 A9 0.28 0.42 2.37 0.83 0.011 0.0030.027 — 0.027 2.83 789 A10 0.18 1.31 1.42 0.35 0.015 0.002 0.052 0.1500.202 3.02 866

TABLE 11 Cooling time (s) Holding time (s) Holding time (s) in theFinishing from completion Cooling in the temperature temperature rangeof Cooling rate (° C./s) Test Steel temperature of finishing rate rangeof 700 to 620° C. within the from 600° C. to Cooling rate (° C./h) No.species (° C.) rolling to 720° C. (° C./s) 720 to 600° C. left-mentionedrange 450° C. below 450° C. 1 A1 854 0.15  900 9.3 6.8 60 22 2 A2 9000.16 1125 4.7 4.4 75 26 3 A3 861 0.13 1090 4.7 3.1 82 24 4 A4 908 0.151253 3.8 3.5 63 24 5 A5 892 0.18  956 3.7 3.4 80 26 6 A6 913 0.18 10722.7 2.4 85 26 7 A7 866 0.13 1123 4.0 3.7 71 27 8 A8 847 0.11 1155 3.33.0 55 26 9 A9 853 0.12 1108 5.3 2.8 59 22 10 A10 872 0.12 1267 3.1 2.754 26 11 A1 868 0.13 1138 1.5 1.3 77 22 12 A4 912 1.25  154 4.1 3.8 6324 13 A8 766 0.07  657 3.5 3.2 62 23 Note: Each underline indicates thatthe relevant value is out of scope of the producing conditions of thepresent invention.

The ferrite crystal grain diameter and grain diameter distribution weredetermined by carrying out crystal orientation analysis at the depth of¼ of the sheet thickness from the sheet surface using the EBSP (electronback scattering pattern) method. Observing the microstructure at thedepth of ¼ of the sheet thickness from the sheet surface as etched withnital or picric acid using a SEM determined the volume fraction of eachphase. The ferrite volume fraction was measured at the depth of ¼ of thesheet thickness from the sheet surface by the so-called mesh method andwas expressed in terms of arithmetic mean. The retained austenite volumefraction was determined by X-ray diffraction measurement. Further, JISNo. 5 test specimens were taken from each rolled material and themechanical characteristics were evaluated on a tensile tester atordinary temperature. As for the thermal stability, each specimen wasimmersed in a salt bath at 700° C. for 10, 30 or 60 minutes and thenrapidly cooled, and the increase rate X (μm/min) in average crystalgrain diameter was calculated by the same method as described above.

For each of the steel sheets according to the present invention, thecrystal grain diameter at the depth of 100 μm from the steel sheetsurface and the crystal grain diameter in the center of the sheetthickness were measured in the same manner as mentioned above. As aresult, in all the steel sheets according to the present invention, thecrystal grain diameter at the depth of 100 μm was not greater than 60%of the grain diameter in the center of the sheet thickness, and thegrain diameter at the depth of ¼ of the sheet thickness was not greaterthan 85% of the grain diameter in the center of the sheet thickness.

These results are shown in Table 12. In Test Nos. 1 to 8 and 10, namelyin the examples according to the present invention, the ferrite occurredas fine grains and the steel sheets were excellent in thermal stabilityand mechanical characteristics. On the contrary, in Test Nos. 9 and 11to 13, namely in the comparatives, the steel sheets were inferior inthermal stability and mechanical characteristics to those obtained inthe examples according to the present invention.

TABLE 12 Increase Average rate X D/3 ≦ Dislocation Retained ferrite(μm/min) d ≦ 3D density Ferrite austenite crystal grain in grain D · XArea (cm⁻²) volume volume Mechanical properties Test Micro- diameter Ddiameter (μm²/ percentage in ferrite fraction fraction TS El TS × El No.structure (μm) at 700° C. min) (%) grains (%) (%) (MPa) (%) (MPa · %) 1α + B + γ 1.91 0.009 0.017 92 6.6 × 10⁷ 77.1 13.7 821 31 25451 Invention2 α + B + γ 2.08 0.010 0.021 89 5.1 × 10⁷ 83.9 10.8 672 33 22176 3 α +B + γ 1.41 0.014 0.020 87 8.8 × 10⁷ 76.2 14.4 861 29 24969 4 α + B + γ2.36 0.010 0.024 94 5.8 × 10⁷ 86.1 9.1 608 34 20672 5 α + B + γ 2.260.007 0.016 92 4.4 × 10⁷ 81.9 11.5 812 32 25984 6 α + B + γ 2.55 0.0090.024 90 4.1 × 10⁷ 82.2 8.6 722 30 21660 7 α + B + γ 2.21 0.009 0.020 916.3 × 10⁷ 83.1 8.9 658 32 21056 8 α + B + γ 1.88 0.010 0.019 91 7.1 ×10⁷ 80.9 11.3 710 33 23430 9 α + B + γ+ cm 1.96 0.012 0.023 81 8.9 × 10⁷38.6 2.1 1265 10 12650 Comparative 10 α + B + γ 2.12 0.030 0.063 71 3.2× 10⁷ 65.5 6.7 785 22 17270 Invention 11 α + B + γ 1.83 0.057 0.104 781.8 × 10⁷ 41.9 8.9 928 17 15776 12 α + B + γ 4.26 0.027 0.115 84 7.6 ×10⁷ 53   6.3 711 20 14220 13 α + B + γ + cm 2.11 0.098 0.207 69 4.6 ×10⁷ 68   2.8 788 11 8668 Note: The symbols in the column ofmicrostructure indicate the following meaning: α: ferrite, B: bainite.γ: retained austenite, cm: granular cementite.

Example 5

Slabs, 50 mm in thickness, made of steel species A to C having therespective chemical compositions shown in Table 13 were hot-rolled incontinuous 6 passes at a total reduction of 96% under the rollingconditions shown in Table 14 and then cooled under the coolingconditions shown in Table 14 to give 2 mm-thick steel sheets.

TABLE 13 Steel species C Si Mn P S sol. Al Cr 2.7 + 5000/(5 + 350 · C +40 · Mn)² Ar₃ (° C.) A 0.06 0.50 2.00 0.009 0.001 0.03 0.2 3.14 790 B0.10 0.03 1.00 0.010 0.003 0.02 — 3.48 820 C 0.20 0.45 2.53 0.008 0.0010.90 — 2.86 850

TABLE 14 Holding Holding time (s) Cooling time time (s) in thetemperature (s) from in the range of 700 completion temperature to 620°C. within Hot-rolled Heating Finishing of finishing Cooling range theCooling rate sheet Chemical temperature temperature rolling to rate of720 to left-mentioned (° C./s) species composition (° C.) (° C.) 720° C.(° C./s) 600° C. range below 600° C. Sample 1 A 1200 830 0.18 800 5 4Cooled to room Invention temperature at 100° C./s Sample 2 B 900 700 — —— 5 Cooled to room Comparative temperature at 100° C./s Sample 3 A 1200830 1.6  100 5 4 Cooled to room temperature at 100° C./s Sample 4 B 1100820 0.14 1200  3 2 Cooled to room Invention temperature at 100° C./sSample 5 B 1200 850 1.5  180 6 5 Cooled to Comparative 500° C./s at 100°C./s and then slowly cooled at 40° C./h Sample 6 C 1150 870 0.12 1000  65 Cooled to Invention 400° C./s at 100° C./s and then slowly cooled at40° C./h Sample 7 C 1000 750 2.0  Air cooling 5 4 Cooled to Comparative400° C./s at 100° C./s and then slowly cooled at 40° C./h Note: Eachunderline indicates that the relevant value is out of scope of theproducing conditions of the present invention.

In the final two passes, the samples 1, 4 and 6 according to the presentinvention and the samples 3 and 5 in the comparatives were rolled at alight reduction of 40 to 35%, and the samples 2 and 7 in thecomparatives were rolled at low temperatures and, in the last pass, itwas rolled at a great reduction of 65%. After rolling, each sample waspickled for scale removal. This sample was cut to pieces with a size of80×200 mm, which were then plated under the conditions described belowusing a vertical hot-dip Zn plating apparatus.

First, each 2.0 mm-thick steel sheet was washed with an NaOH solution at75° C. for degreasing and then annealed at 600, 720 or 840° C. for 60seconds in an atmosphere consisting of atmosphere gas N₂+20% H₂ andhaving a dew point of −40° C. After annealing, the steel sheet wascooled to the vicinity of the bath temperature, dipped in one of severalplating baths for 3 seconds and then adjusted to a plating metalcoverage of 50 g/m² per side by the wiping method. In the case ofalloying treatment, the plated sheet was then subjected to 30 seconds ofheat treatment at 500° C. using an infrared heating apparatus. Thecooling rate was adjusted by varying the gas quantity and mist quantity.The temper rolling after plating was carried out at an arithmetical meanroll roughness Ra of 1 to 5 μm and a load of 200 tons/m.

The ferrite crystal grain diameter and grain diameter distribution atthe depth of ¼ of the sheet thickness from the steel sheet surface afterhot rolling or after plating were examined by the EBSP method, thevolume fraction of each phase in the steel microstructure was examinedby observation under a SEM and by X-ray diffraction measurement of theetched microstructure and, further, the dislocation density in ferritecrystal grains was examined under a transmission electron microscope.Further, for each steel sheet after plating, JIS No. 5 tensile testspecimens were taken and examined for tensile characteristics.

For each of the steel sheets according to the present invention, thecrystal grain diameter at the depth of 100 μm from the steel sheetsurface and the crystal grain diameter in the center of the sheetthickness were measured in the same manner as mentioned above. As aresult, in all the steel sheets according to the present invention, thecrystal grain diameter at the depth of 100 μm was not greater than 60%of the grain diameter in the center of the sheet thickness, and thegrain diameter at the depth of ¼ of the sheet thickness was not greaterthan 85% of the grain diameter in the center of the sheet thickness.

These results are shown in Table 15. Here, in each of Test Nos. 1 to 3,6 to 8, and 12, the hot-rolled steel sheet having a fine ferrite crystalgrain microstructure according to the present invention had a high levelof thermal microstructure stability and, therefore, even after hot-dipplating treatment, the ferrite crystal grain diameter showed almost noincrease and the fine ferritic microstructure retained an appropriategrain diameter distribution and a low dislocation density. Therefore,after plating as well, the steel sheet was excellent in both mechanicalcharacteristics and thermal stability. On the contrary, in thecomparatives, namely in Test Nos. 4, 5, 9 and 11, the thermal stabilityand mechanical characteristics were inferior as compared with those inthe examples according to the present invention.

TABLE 15 Average ferrite Ferrite D/3 ≦ d ≦ 3D crystal grain volume AreaHot rolled Alloying diameter D fraction percentage sheet species SoakingPlating bath treatment (μm) (%) (%) Test (Sample temperature Platingtemperature temperature After hot After After hot After After hot AfterNo. No.) (° C.) bath (° C.) (° C.) rolling plating rolling platingrolling plating 1 1 600 Zn—0.2% Al 460 — 2.2 2.2 89 89 97 97 2 1 720Zn—0.2% Al 460 — ″ 2.2 ″ 89 ″ 98 3 1 840 Zn—0.2% Al 460 — ″ 2.6 ″ 85 ″98 4 2 840 Zn—0.2% Al 460 — 3.0 4.2 92 84 77 65 5 3 840 Zn—0.2% Al 460 —6.2 6.5 85 84 93 95 6 4 840 Zn—5% Al 420 — 2.4 2.7 90 86 95 97 7 4 840Zn—5% Al—3% 420 — ″ 2.6 ″ 87 ″ 96 Mg 8 4 840 Zn—55% Al—1.6% 600 — ″ 2.7″ 88 ″ 94 Si 9 5 840 Zn—5% Al—1.6% 600 — 5.4 6.8 92 87 92 90 Si 10  6840 Zn—0.1Al % 460 500 1.8 2.1 75 74 93 95 11  7 840 Zn—0.1% Al 460 5002.6 4.5 60 75 77 90 12  6 840 Al 10% Si 620 — 1.8 2.1 75 74 93 94Dislocation Second phase species density (cm-2) and volume fraction D ·X in ferrite grains (%) (μm²/min) Mechanical properties after platingTest After hot After After hot After After hot After YP TS El YR TS × ElNo. rolling plating rolling plating rolling plating (MPa) (MPa) (%) (—)(MPa · %) 1 7.6 × 10⁷ 7.2 × 10⁷ M(11) tM(11) 0.025 0.025 Invention 2 ″5.6 × 10⁷ ″ ″ 0.025 0.029 3 ″ 4.6 × 10⁷ ″ M(15) 0.025 0.032 495 820 22.10.60 18122 4  2.3 × 10¹⁰ 3.1 × 10⁷ M(8) M(16) 0.102 0.110 420 650 24.60.65 15990 Comparative 5 5.6 × 10⁷ 3.9 × 10⁷ M(15) M(16) 0.062 0.053 410645 25.2 0.64 16254 6 8.2 × 10⁷ 4.2 × 10⁷ B(8), M(1), B(9), P(5) 0.0190.026 525 560 29.5 0.94 16520 Invention cm(1) 7 ″ 3.7 × 10⁷ B(8), M(1),B(10), P(5) 0.019 0.022 512 565 28.3 0.91 15990 cm(1) 8 ″ 4.5 × 10⁷B(8), M(1), B(8), P(6) 0.019 0.021 488 542 31.0 0.90 16802 cm(1) 9 2.9 ×10⁸ 7.7 × 10⁷ P(8) B(7), P(6) 0.051 0.045 333 445 34.8 0.75 15486Comparative 10  9.8 × 10⁷ 3.5 × 10⁷ γ(18), B(7) γ(20), B(6) 0.019 0.031688 810 29.6 0.85 23976 Invention 11   4.1 × 10¹⁰ 8.6 × 10⁷ γ(8), B(32)γ(15), B(10) 0.132 0.156 593 760 23.2 0.78 17632 Comparative 12  7.5 ×10⁷ 4.1 × 10⁷ γ(18), B(7) γ(15), B(8) 0.019 0.025 707 852 23.0 0.8319596 Invention Note: The symbols in the column of microstructureindicate the following meaning: M: martensite, tM: tempered martensite,B: bainite. γ: retained austenite, cm: granular cementite.

Example 6

The 2 mm-thick steel sheets obtained in Examples 1, 3 and 5 weresubjected to penetration butt welding using the plasma welding method(welding speed: 0.5 m/min, welding current: about 180 A) and the laserwelding method (welding speed: 1.0 m/min, concentration spot diameter:0.6 mm, output: 3000 W). The main composition of each steel sheet andthe carbon equivalent Ceq(I) are shown in Table 16.

TABLE 16 Test Main chemical composition No. Steel sheet sample C Si MnCr Mo Ni V Ceq(I) (%) 1 Example 1, Test No. 2 0.105 0.05 1.03 0.279 2 3Example 1, Test No. 9 0.075 0.03 0.44 0.150 4 Example 1, Test No. 200.080 0.06 0.70 0.31 0.3 0.268 5 Example 1, Test No. 24 0.045 0.42 0.700.1 0.186 6 Example 3, Test No. A1 0.050 0.03 2.49 0.466 7 Example 3,Test No. A4 0.060 0.49 1.98 0.26 0.462 8 Example 3, Test No. A7 0.0500.53 1.56 0.24 0.392 9 Example 5, Test No. 1 0.060 0.50 2.00 0.20 0.45410 11 Example 5, Test No. 6 0.100 0.03 1.00 0.268 12 Example 1, Test No.10 0.030 0.03 0.22 0.068 13 Example 1, Test No. 25 0.023 0.02 0.18 0.05414 Example 5, Test No. 10 0.200 0.45 2.53 0.640 15 Example 5, Sample 2,no plating treatment 0.060 0.50 2.00 0.20 0.454 Note: Each underline inthe Ceq(I) column indicates that the relevant value is out of scope ofthe preferred mode of embodiment of the present invention.

The characteristics of each weld obtained were evaluated by punchstretch forming testing using a ball head with a diameter of 50 mm. Theshape of the ball head punch stretch forming test specimen is shown inFIG. 1. The state in which the direction of principal strain is parallelto the direction of the weld line 2 is shown in FIG. 1 (a) (Type I) andthe state in which the direction of principal strain is perpendicular tothe direction of the weld line 2 is shown in FIG. 1 (b) (Type II). Ballhead punch stretch forming test specimen 1 were excised from each weldand evaluated for stretch height and site of rupture.

The results are shown in Table 17. Here, in Test No. 13, the steel sheethad a low Ceq(I) value and the bead melted and solidified on theoccasion of welding was soft as compared with the base metal andunderwent rupture in Type II testing. In Test No. 14, the steel sheethad an excessively great Ceq(I) value, so that the bead hardened to anexcessive extent and underwent bead rupture in type I testing. In TestNo. 15, the steel sheet was a fine-grained microstructure steel sheetproduced by low-temperature rolling and, therefore, was inferior inthermal stability and rupture occurred in the HAZ site in type IItesting. On the contrary, in Test Nos. 1 to 12, it is seen that thesteel sheets still showed high workability in a weld-including regionafter fusion welding using plasma or laser beams and thus were excellentin post-welding formability.

TABLE 17 Welding conditions TYPE I (parallel) TYPE II (vertical) WeldingStretch Stretch Welding rate height Site of height Site of Test No.method (m/min) (mm) rupture (mm) rupture 1 Plasma 0.5 17.0 Base metal17.5 Base metal 2 Laser 1.0 21.0 Base metal 21.5 Base metal 3 Plasma 0.520.0 Base metal 23.0 Base metal 4 Plasma 0.5 18.0 Base metal 17.5 Basemetal 5 Plasma 0.5 18.0 Base metal 18.0 Base metal 6 Plasma 0.5 16.0Base metal 16.5 Base metal 7 Plasma 0.5 17.0 Base metal 16.5 Base metal8 Plasma 0.5 17.0 Base metal 16.5 Base metal 9 Plasma 0.5 17.5 Basemetal 16.5 Base metal 10 Laser 1.0 19.5 Base metal 18.5 Base metal 11Plasma 0.5 18.0 Base metal 18.5 Base metal 12 Plasma 0.5 18.5 Base metal18.6 Base metal 13 Plasma 0.5 18.5 Base metal 16.5 Bead cracking 14Plasma 0.5 10.0 Bead cracking 16.0 Base metal 15 Plasma 0.5 16.0 Basemetal 10.5 HAZ site

Example 7

Those 2 mm-thick fine-grained hot-rolled steel sheets which wereobtained in Examples 1, 3, 4 and 5 and had a 440 to 780 MPa class ofstrength in terms of tensile strength TS as well as commerciallyavailable coarse-grained hot-rolled steel sheets almost parallel intensile strength to those just mentioned above were used and evaluatedfor resistance weldability. The tensile strength TS, the main chemicalcomposition and the carbon equivalent Ceq(II) of each of thosehot-rolled steel sheets are shown, together with other data, in Table18.

TABLE 18 Average ferrite 2.7 + 5000/ Test TS Main chemical compositionCeq(II) crystal grain (5 + 350 · D · X No. Steel sheet sample (MPa) C SiAl Mn Cr (%) Rsp diameter D (μm) C + 40 · Mn)² (μm²/min) 1 Example 1,Test No. 25 423 0.023 0.02 0.032 0.18 — 0.025 13.9 6.50 14.89 0.078 2Example 1, Test No. 10 434 0.030 0.03 0.030 0.22 — 0.033 14.2 4.25 11.170.021 3 Commercial hot-rolled steel 440 0.130 0.05 0.020 1.05 — 0.14118.8 7.30 3.28 0.068 sheet 4 Example 1, Test No. 2 578 0.105 0.05 0.0301.03 — 0.116 18.8 2.38 3.43 0.024 5 Example 1, Test No. 17 600 0.1550.03 0.030 0.62 — 0.162 16.4 2.95 3.41 0.023 6 Commercial hot-rolledsteel 590 0.060 0.60 0.034 1.60 0.21 0.085 30.5 6.50 3.32 0.069 sheet 7Example 3, Test No. A1 813 0.050 0.03 0.034 2.49 — 0.078 27.9 2.64 3.040.026 8 Example 3, Test No. A4 788 0.060 0.49 0.030 1.98 0.26 0.087 31.02.83 3.15 0.022 9 Example 5, Test No. 3 820 0.060 0.50 0.030 2.00 0.200.088 31.2 2.60 3.14 0.032 10 Example 5, Test No. 10 810 0.200 0.450.900 2.53 — 0.230 44.1 2.10 2.86 0.031 11 Example 4, Test No. 5 8120.110 1.75 0.360 1.22 — 0.142 48.9 2.26 3.29 0.016 12 Example 4, TestNo. 10 785 0.180 1.31 0.350 1.42 — 0.209 42.3 2.12 3.02 0.063 13Commercial hot-rolled steel 785 0.098 0.74 0.025 1.80 0.10 0.125 32.85.20 3.10 0.086 sheet 14 Example 4, Test No. 13 788 0.140 1.48 0.0201.66 — 0.173 41.4 2.11 3.04 0.207 Note: Each underline in the Ceq(II)column indicates that the relevant value is out of scope of thepreferred mode of embodiment of the present invention.

Test pieces, 30×100 mm in size, were excised and each pair of pieceswere placed one on top of the other with a lap margin of 30 mm and ajoint was formed under the conditions of a pressure of 3920 N (400 kg·f)and 30 cycles of electric current application using dome-shapedelectrodes with a diameter of 8 mm while the welding current was varied.

The expulsion-causing current was measured, and the joint was subjectedto shear tensile testing to evaluate the same for maximum load atrupture. Further, the cross section of the spot weld was observedmacroscopically and the nugget diameter after picric acid etching wasmeasured and thereby the electric current range from the formation of anugget diameter of 4√{square root over ( )}t to the occurrence ofexpulsion and the current range from the button breakage of the joint tothe occurrence of expulsion were deter-mined for resistance weldabilityevaluation. The maximum load at rupture for the joint on which thesmallest button diameter was obtained among the button break-age-causingconditions was taken as the maximum load at joint rupture.

The results are shown in Table 19. Here, in Test No. 1, the Ceq(II) wasinsufficient, so that the maximum hardness of the weld was low and thejoint strength was also low. In Test Nos. 10 to 12, the Ceq(II) or Rspvalue was excessively high, so that, in each case, the current rangefrom button breakage to expulsion was narrow. In Test No. 14, the steelsheet was a fine-grained microstructure steel sheet produced bylow-temperature rolling and was inferior in thermal stability and,therefore, in spite of the fact that the bead portion had an appropriatehardness, the HAZ site became softened and, therefore, the jointstrength was low. In Test Nos. 3, 6 and 13, the steel sheets used werecommercial coarse-grained hot-rolled steel sheets and were inferior inthermal stability and, moreover, the current range from button breakageto expulsion was narrow for all of them. On the contrary, the steelsheets used in Test Nos. 2, 4, 5 and 7 to 9 had excellent mechanicalcharacteristics and at the same time had a broad current range properfor welding and showed good resistance weldability.

TABLE 19 Proper current range (kA) Current range from 4√ T Current rangeMaximum joint nugget diameter from button Test rupture load formation tobreakage to Strength No (kN) expulsion expulsion level 1 16.0 5.1 3.6440 MPa 2 20.0 4.8 3.5 class 3 20.0 4.5 2.8 4 25.0 4.0 2.5 590 MPa 525.5 3.0 1.5 class 6 25.5 3.0 1.0 7 35.5 3.8 2.2 780 MPa 8 35.0 3.0 1.8class 9 35.0 2.8 1.8 10 35.0 2.5 0.3 11 35.0 2.4 0.4 12 35.0 2.7 0.4 1335.0 2.8 0.4 14 26.0 2.8 0.4

Example 8

The steel having the chemical composition Al as shown in Table 1 inExample 1 was prepared by melting and worked to a thickness of 30 mm byhot forging. Thereafter, the semi-finished product was reheated to 1000°C. and then rolled on a small-sized test tandem mill to give a finished1.3 mm-thick sheet. The rolling finish temperature was 830° C., whichwas higher than the Ar₃ point. The final three rolling passes, a highreduction of 40 to 50% was employed. After rolling finish, water coolingwas started after the lapse of 0.05 second and the sheet was cooled to atemperature of 680° C. at a cooling rate of not lower than 1000° C./secand thereafter allowed to cool for about 4 seconds and, after arrival ofthe temperature at 600° C., again cooled with water until it was cooledto room temperature. Thereafter, it was pickled, cold-rolled to athickness of 0.5 mm at a reduction of 62% and then annealed. Theannealing was carried out by dipping in a salt bath at 800° C. for about2 minutes, followed by water cooling to room temperature.

After hot rolling, the average ferrite crystal grain diameters at thedepths of 1/16 and ¼ of the sheet thickness from the surface and at thecenter of the sheet thickness were 1.4, 1.7 and 2.3 μm, respectively.The ferrite volume fraction was 93% and the second phase was bainite ormartensite. At the depth of ¼ of the sheet thickness from the surface,90% or more of ferrite grains had a grain diameter falling within therange from ⅓ of to 3 times the average grain diameter at that depth, andthe product D·X of the increase rate in grain diameter at 700° C. andthe grain diameter was 0.026 μm²/min (=0.015 μm/min×1.7 μm). FIG. 2shows the time course of the changing ferrite grain diameter at 700° C.This figure indicates that the grain growth rate was as low as 0.015μm/min. The ferrite grains had an equiaxed shape.

The changes in grain diameter after cold rolling and annealing dependingon the annealing time are shown in FIG. 3. The data at annealing time 0are the grain diameter data after hot rolling. Although the graindiameters increased as a result of annealing, the changes were small;even after annealing, the grain diameters at the depths of 1/16 and ¼ ofthe sheet thickness from the surface and at the center of the sheetthickness were 2.2, 2.4 and 2.9 μm, respectively, and the grains werestill fine. No grain coarsening occurred due to the prolongation of theannealing time. A commercial cold-rolled steel sheet having the samecomposition had a crystal grain diameter of about 7.5 μm, while thegrain diameter of the steel sheet according to the present invention wasabout ⅓ of that of the commercial sheet. An example of themicrostructure on the occasion mentioned above (5 minutes of annealingat 800° C.) is shown in FIG. 4. The ferrite volume fraction of the steelsheet according to the present invention after annealing was about 70%and the second phase was a martensite. The product D·X of the increaserate in grain diameter at 700° C. and the grain diameter was about 0.02μm²/min. At the depth of ¼ of the sheet thickness from the surface, 90%or more of ferrite grains had a grain diameter falling within the rangefrom ⅓ of to 3 times the average grain diameter at that depth.

Example 9

The hot-rolled steel sheet of Example 8 was likewise pickled, thencold-rolled to a sheet thickness of 0.5 mm at a reduction of 62%, andannealed. In the annealing, a heat treatment step in an industrialcontinuous annealing line was simulated. The rate of temperatureincrease was 10 to 15° C./sec, the soaking temperature was 750° C. or800° C., and the post-rolling cooling conditions were equivalent to thecontinuous Zn alloy plating treatment conditions when the soakingtemperature was 750° C. or, when the soaking temperature was 800° C.,overaging treatment comprising slow cooling from 400° C. to 320° C. wasadded.

The ferrite grain diameters after cold rolling and heat treatment were3.5, 3.8 and 4.1 μm at the depths of 1/16 and ¼ of the sheet thicknessfrom the surface and at the center of the sheet thickness, respectively,when the soaking temperature was 750° C. and, when the soakingtemperature was 800° C., they were 4.2, 4.6 and 5.0 μm, respectively.These grain diameters are 50 to 60% of the crystal grain diameter of acommercial cold-rolled steel sheet having the same composition, namelyabout 7.5 μm. An example of the microstructure of the material obtainedby using the soaking temperature of 750° C. is shown in FIG. 5. Theproduct D·X of the increase rate in grain diameter at 700° C. and thegrain diameter was not greater than 0.01 μm²/min, and the grain diameterscarcely changed during the measurement period (30 minutes). At thedepth of ¼ of the sheet thickness from the surface, 90% or more offerrite grains had a grain diameter falling within the range from ⅓ ofto 3 times the average grain diameter at that depth. In each steelsheet, the ferrite volume fraction was not lower than 93%, and thesecond phase was pearlite. The mechanical characteristics of these steelsheets are shown in Table 20. The data indicate that the yield strengthsof the steel sheets of the invention were higher by 60 to 80 MPa and thetensile strengths were also higher by 30 to 50 MPa than those of thecommercial steel sheet identical in composition. They were almostparallel in uniform elongation (UEL) to the commercial steel sheet inspite of the increases in strength. The total elongation EL values werelower but this is because the commercial steel sheet was as thick as 1.2mm. Considering the difference in sheet thickness, the steel sheets ofthe invention are parallel or superior in strength-elongation balance tothe commercial steel sheet.

TABLE 20 Material YS (MPa) TS (MPa) UEL (%) EL (%) YR Invention 800° C.422.9 493.3 20.6 31.0 0.86 Invention 750° C. 412.2 472.4 22.7 33.4 0.87Commercial steel 345 440 23 40.5 0.78

It is seen that the yield strengths of the steel sheets of the inventionwere higher by 60 to 80 MPa and the tensile strengths were also higherby 30 to 50 MPa than those of the commercial steel sheet identical incomposition. They were almost parallel in uniform elongation (UEL) tothe commercial steel sheet in spite of the increases in strength. Thetotal elongation EL values were lower but this is because the commercialsteel sheet was as thick as 1.2 mm. Considering the difference in sheetthickness, the steel sheets of the invention are parallel or superior instrength-elongation balance to the commercial steel sheet.

INDUSTRIAL APPLICABILITY

The steel sheet according to the present invention has ultra finecrystal grains and is excellent in thermal stability, hence can endurethe heat during welding or plating, and also is excellent in mechanicalcharacteristics. Such steel sheet excellent in thermal stability andmechanical characteristics can be produced with ease by the methodaccording to the present invention.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 This figure shows the shape of a ball head punch stretch formingtest specimen. FIG. 1 (a) shows the state in which the direction ofprincipal strain is parallel to the direction of the weld line, and FIG.1 (b) shows the state in which the direction of principal strain isperpendicular to the direction of the weld line.

FIG. 2 This figure shows the time course of the changing ferrite graindiameter at the depth of ¼ of the sheet thickness from the surface.

FIG. 3 This figure shows the changes, depending on the annealing time,of the ferrite grain diameter after cold rolling and annealing.

FIG. 4 This figure shows the microstructure after 5 minutes of annealingat 800° C. following cold rolling.

FIG. 5 This figure shows an example of the microstructure afterannealing at 750° C. following cold rolling.

EXPLANATION OF SYMBOLS

-   1—Ball head punch stretch forming test specimen-   2—Weld line

1. A hot-rolled steel sheet of carbon steel or low-alloy steel, the mainphase of which is ferrite, and is characterized in that the averageferrite crystal grain diameter D (μm) at the depth of ¼ of the sheetthickness from the steel sheet surface satisfies the relationsrespectively defined by the formulas (1) and (2) given below and theincrease rate X (μm/min) in average ferrite crystal grain diameter at700° C. at the depth of ¼ of the sheet thickness from the steel sheetsurface and said average crystal grain diameter D (μm) satisfy therelation defined by the formula (3) given below:1.2≦D≦7  formula (1)D≦2.7+5000/(5+350.C+40.Mn)²  formula (2)D·X≦0.1  formula (3) wherein C and Mn represent the contents (in % bymass) of the respective elements in the steel.
 2. A hot-rolled steelsheet according to claim 1, characterized in that, at the depth of ¼ ofthe sheet thickness from the steel sheet surface, the area percentage offerrite crystal grains the crystal grain diameter d (μm) of whichsatisfies the relation defined by the formula (4) given below is atleast 80%:D/3≦d≦3D  formula (4) wherein D represents the average ferrite crystalgrain diameter (μm) at the depth of ¼ of the sheet thickness from thesteel sheet surface.
 3. A cold-rolled steel sheet of carbon steel orlow-alloy steel, the main phase of which is ferrite, and ischaracterized in that the average ferrite crystal grain diameter D (μm)at the depth of ¼ of the sheet thickness from the steel sheet surfacesatisfies the relations respectively defined by the formulas (5) and (6)given below and the increase rate X (μm/min) in average ferrite crystalgrain diameter at 700° C. at the depth of ¼ of the sheet thickness fromthe steel sheet surface and said average crystal grain diameter D (μm)satisfy the relation defined by the formula (3) given below:1.2≦D≦9.3  formula (5)D≦5.0−2.0.Cr+5000/(5+350.C+40.Mn)²  formula (6)D·X≦0.1  formula (3) and, further, that, at the depth of ¼ of the sheetthickness from the steel sheet surface, the area percentage of ferritecrystal grains the crystal grain diameter d (μm) of which satisfies therelation defined by the formula (4) given below is at least 80%:D/3≦d≦3D  formula (4) wherein C, Cr and Mn represent the contents (in %by mass) of the respective elements in the steel.
 4. A hot-rolled orcold-rolled steel sheet according to any of claims 1 to 3, characterizedin that it contains, as a second phase other than ferrite, a total ofless than 50% of one or more species selected from the group consistingof less than 50% of bainite, less than 30% of pearlite, less than 5% ofgranular cementite, less than 5% of martensite and less than 3% ofretained austenite and has a yield ratio lower than 0.75, % in eachoccurrence being % by volume.
 5. A hot-rolled or cold-rolled steel sheetaccording to any of claims 1 to 3, characterized in that it contains, asa second phase other than ferrite, 5 to 40% by volume of martensite andhas a yield ratio lower than 0.75.
 6. A hot-rolled or cold-rolled steelsheet according to any of claims 1 to 3, characterized in that itcontains, as a second phase other than ferrite, 3 to 30% by volume ofretained austenite and has a product, TS×El, of tensile strength TS(MPa) and total elongation El (%) of not less than 18000 (MPa·%).
 7. Ahot-rolled steel sheet according to any of claims 1, 2, 4, 5 and 6,characterized in that the average crystal grain diameter Ds (μm) at thedepth of 1/16 of the sheet thickness from the steel sheet surface, theaverage crystal grain diameter D (μm) at the depth of ¼ of the sheetthickness from the steel sheet surface and the average crystal graindiameter Dc (μm) at the center of the sheet thickness satisfy therelations Ds≦0.75D and D≦0.9Dc.
 8. A cold-rolled steel sheet accordingto any of claims 3 to 6, characterized in that the average crystal graindiameter Ds (μm) at the depth of 1/16 of the sheet thickness from thesteel sheet surface and the average crystal grain diameter Dc (μm) atthe center of the sheet thickness satisfy the relation Ds≦0.9Dc.
 9. Ahot-rolled or cold-rolled steel sheet according to any of claims 1 to 8,characterized in that the carbon equivalent Ceq(I) defined by theformula (7) given below is 0.06 to 0.6%:Ceq(I)=C+Mn/6+Si/24+Cr/5+Mo/4+Ni/40+V/14  formula (7) wherein thesymbols of elements in the above formula represent the contents (in % bymass) of the respective elements in the steel.
 10. A hot-rolled orcold-rolled steel sheet according to any of 1 to 8, characterized inthat the C content is not higher than 0.17% by mass, the carbonequivalent Ceq(II) defined by the formula (8) given below is 0.03 to0.20% and the base metal resistance indicator Rsp defined by the formula(9) given below is not higher than 45:Ceq(II)=C+Mn/100+Si/90+Cr/100  formula (8)Rsp=13.5×(Si+Al+0.4Mn+0.4Cr)+12.2  formula (9) wherein the symbols ofelements in the above formulas represent the contents (in % by mass) ofthe respective elements in the steel.
 11. A hot-dip-plated hot-rolled orcold-rolled steel sheet characterized in that it comprises a Zn, Al,Zn—Al alloy or Fe—Zn alloy coat layer formed on the surface of ahot-rolled steel sheet according to any of claims 1 to
 10. 12. A methodof producing a hot-rolled steel sheet according to any of claims 1, 2,4, 5, 6, 7, 9, 10 and 11 by subjecting a carbon steel or low-alloy steelslab to multi-pass hot rolling to give the hot-rolled steel sheet,characterized in that the final rolling pass is finished at atemperature not lower than the Ar₃ point and not lower than 780° C. andthen the rolled sheet is cooled to 720° C. or below within 0.4 second ata cooling rate of not lower than 400° C./second and then maintained in atemperature range of 600 to 720° C. for at least 2 seconds.
 13. A methodof producing a cold-rolled steel sheet according to any of claims 3, 4,5, 6, 8, 9 and 10, characterized in that a hot-rolled steel sheetobtained by the method defined in claim 12 is pickled, then cold-rolledat a reduction of 40 to 90% and thereafter heat-treated at a temperatureof not higher than 900° C.
 14. A method of producing a hot-dip-platedhot-rolled or cold-rolled steel sheet according to claim 11,characterized in that a hot-rolled steel sheet obtained by the methoddefined in claim 12 is subjected to pickling or to picking and furthercold rolling at a reduction of 40 to 90%, and then to hot-dip plating ina continuous hot-dip plating line.